i
Synthesis, Characterisation and Deposition of Nano
Hydroxyapatite coatings on Bio-degradable
Magnesium for Potential Orthopaedic Applications
Sridevi Brundavanam, M.Sc
This thesis is presented for the degree of Doctor of Philosophy of Murdoch University
2015
ii
Declaration
I declare that this thesis is my own account of my research and contains as
its main content, work which has not previously been submitted for a
degree at any tertiary education institution.
.....................................................................
Sridevi Brundavanam
iii
Dedicated to my daughter
Gayatri
&
My husband
Ravi Krishna
iv
Abstract
Today, Magnesium (Mg) based alloys are receiving increasing attention as potential
biodegradable implant materials for orthopaedic applications. Despite advantageous
properties such as density and elastic modulus that are similar to bone, magnesium’s
rapid degradation rate when immersed in the highly corrosive body fluid environment
has severely limited its clinical application. The focus of this thesis research was to
develop biocompatible calcium phosphate coatings with tuneable biodegradable
properties that are capable of extending the operational life of an Mg based implant and
allow sufficient time for tissue regeneration to take place.
The research developed and examined three types of calcium phosphate coatings
designed to reduce the degradation rate and prolong the life of Mg test substrates.
Dicalcium phosphate dihydrate (DCPD) or Brushite coatings were formed on Mg
substrates via a straightforward chemical immersion technique. While amorphous
calcium phosphate (ACP) coatings were formed on Mg substrates using an
electrochemical technique. Brushite coatings were characterised by widespread flower-
like surface structures and the ACP coating were granular in structure with a surface
covering of tube-like structures. The third coating examined was formed by the
subsequent transformation of Brushite and ACP coatings to hydroxyapatite (HAP) via a
low-temperature hydrothermal process. Importantly, HAP is the mineral component
found in bone and is known to promote bone cell adhesion, differentiation and
osteointegration.
v
Advanced characterisation techniques such as X-ray diffraction, field emission scanning
electron microscopy, energy dispersive spectroscopy and Fourier transform infrared
spectroscopy were used to investigate the size, morphology, crystalline structure,
composition and topographical features of both uncoated and coated substrates.
Degradation behaviour studies were carried on pure Mg substrates and the various
substrate coatings using two simulated body fluid electrolytes at human body
temperature (37 ºC). The first was phosphate buffer saline (PBS) solution and the
second was Ringer’s solution. Corrosion rate measurements revealed DCPD coated
substrates had the lowest corrosion rate in both PBS (0.126 mm/yr) and Ringer’s
solution (0.1 mm/yr) compared to HAP coated [PBS (0.279 mm/yr) and Ringer’s
solution (0.264 mm/yr)] and uncoated Mg substrates [PBS (1.829 mm/yr) and Ringer’s
solution (3.828 mm/yr)]. The results of the research indicate that both DCPD and HAP
coatings have the potential to reduce the corrosive effects produced by both PBS and
Ringer’s solution. The results also suggest that the coatings have the ability to reduce
the degradation rate of Mg substrates in the physiological environment.
Furthermore, because of HAPs complex hexagonal structure it offers an effective high
capacity absorbent matrix and the present research has shown that like bone HAP can
accumulate metallic materials such as cadmium, copper, iron and zinc. The significant
improvement in corrosion resistance and the ability of HAP coatings in particular to
behave as a temporary repository of metallic materials during degradation is an
important step in the development of a biodegradable Mg implant for orthopaedic
applications.
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Publications
1. Gérrard Eddy Jai Poinern, Sridevi Brundavanam and Derek Fawcett, Biomedical
Magnesium Alloys: A Review of Material Properties, Surface Modifications and
its Potential as a Biodegradable Orthopaedic Implant. American Journal of
Biomedical Engineering. 2012; 2 (6); 218-240.
2. Sridevi Brundavanam, Gérrard Eddy Jai Poinern, Derek Fawcett, Chemical
immersion coatings to improve biological degradability of magnesium
substrates for potential orthopaedic applications. International Journal of
Biomedical Materials Research. 2014; 2(2): 7-14.
3. Sridevi Brundavanam, Gérrard Eddy Jai Poinern, Derek Fawcett, Synthesis of a
Hydroxyapatite nanopowder via ultrasound irradiation from calcium hydroxide
powders for potential biomedical applications. Nanoscience and
Nanoengineering. 2015; 3(1): 1-7.
4. Sridevi Brundavanam, Gérrard Eddy Jai Poinern, Derek Fawcett. Kinetic and
adsorption behaviour of aqueous Fe (II), Cu (II) and Zn (II) using a 30
nanometre scale hydroxyapatite powder synthesized via a combined ultrasound
and microwave based technique. American Journal of Materials Science 2015;
5(2): 31-40.
5. Sridevi Brundavanam, Gérrard Eddy Jai Poinern, Derek Fawcett. Growth of
flower-like brushite structures on magnesium substrates and their subsequent
low temperature transformation to hydroxyapatite. American Journal of
Biomedical Engineering. 2014; 4(4): 79-87.
6. Sridevi Brundavanam, Gérrard Eddy Jai Poinern, Derek Fawcett.
Electrochemical synthesis of micrometre amorphous calcium phosphate tubes
and their transformation to hydroxyapatite tubes. International Journal of
Sciences. 2015; 4(2):7-15.
7. Derek Fawcett, Sridevi Brundavanam, Gérrard Eddy Jai Poinern. Growth and
corrosion Behaviour of amorphous micrometre scale calcium phosphate
coatings on Magnesium substrates. International Journal of Materials
Engineering 2015; 5(1): 10-16.
8. Gérrard Eddy Jai Poinern, Sridevi Brundavanam, Derek Fawcett. Kinetic and
adsorption behaviour of aqueous cadmium using a 30 nanometre scale
hydroxyapatite powder synthesized via a combined ultrasound and microwave
based technique (Submitted to Journal of Materials and Environmental Science)
vii
Table of Contents
Page
DECLARATION ii
DEDICATION iii
ABSTRACT iv
PUBLICATIONS vi
TABLE OF CONTENTS vii
ACKNOWLEDGMENTS xi
Chapter 1 – Introduction 1
Chapter 2 – Literature Review
2.1. Overview and author contributions
2.2. Publications:
Review Article 1
Gérrard Eddy Jai Poinern, Sridevi Brundavanam and Derek Fawcett,
Biomedical Magnesium Alloys: A Review of Material Properties, Surface
Modifications and its Potential as a Biodegradable Orthopaedic Implant.
American Journal of Biomedical Engineering. 2012; 2 (6); 218-240.
Review Article 2
Sridevi Brundavanam, Gérrard Eddy Jai Poinern, Derek Fawcett, Chemical
immersion coatings to improve biological degradability of magnesium
substrates for potential orthopaedic applications. International Journal of
Biomedical Materials Research. 2014; 2(2): 7-14.
2.3. Chapter Summary
11
11
12
13
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Chapter 3 – Case Study 1: Synthesis of Hydroxyapatite
Nanometre Scale Powders
3.1. Overview and author contributions
3.2. Optimising Laboratory Procedure for Synthesizing Hydroxyapatite via a
combined Ultrasound and Microwave based technique.
3.3. Published Research Article:
Case Study 1
Sridevi Brundavanam, Gérrard Eddy Jai Poinern, Derek Fawcett, Synthesis
of a Hydroxyapatite nanopowder via ultrasound irradiation from calcium
hydroxide powders for potential biomedical applications. Nanoscience and
Nanoengineering. 2015; 3(1): 1-7.
3.4. Chapter Summary
14
14
16
19
Chapter 4 – Case Studies 2 and 3: Adsorption of Selected Heavy
metals using Hydroxyapatite absorbers
4.1. Overview and author contributions
4.2. Published Research Articles
Case Study 2
Gérrard Eddy Jai Poinern, Sridevi Brundavanam, Derek Fawcett. Kinetic
and adsorption behaviour of aqueous cadmium using a 30 nanometre scale
hydroxyapatite powder synthesized via a combined ultrasound and
microwave based technique (Currently under review - submitted to Journal
of Materials and Environmental Science)
Case Study 3
Sridevi Brundavanam, Gérrard Eddy Jai Poinern, Derek Fawcett. Kinetic
and adsorption behaviour of aqueous Fe (II), Cu (II) and Zn (II) using a 30
21
21
23
ix
nanometre scale hydroxyapatite powder synthesized via a combined
ultrasound and microwave based technique. American Journal of Materials
Science 2015; 5(2): 31-40.
4.3. Chapter Summary
52
Chapter 5 – Case Study 4: Growth of flower-like brushite
structures on magnesium substrates and their subsequent low
temperature transformation to hydroxyapatite.
5.1. Overview and author contributions
5.2. Published Research Article:
Case Study 4
Sridevi Brundavanam, Gérrard Eddy Jai Poinern, Derek Fawcett. Growth of
flower-like brushite structures on magnesium substrates and their
subsequent low temperature transformation to hydroxyapatite. American
Journal of Biomedical Engineering. 2014; 4(4): 79-87.
5.3. Chapter Summary
53
53
54
55
Chapter 6 – Case Studies 5 and 6: Electrochemical synthesis of
amorphous calcium phosphate coating on magnesium substrates
in aqueous solutions.
6.1. Overview and author contributions
6.2. Published Research Articles
Case Study 5
Sridevi Brundavanam, Gérrard Eddy Jai Poinern, Derek Fawcett.
Electrochemical synthesis of micrometre amorphous calcium phosphate
56
56
57
x
tubes and their transformation to hydroxyapatite tubes. International
Journal of Sciences. 2015; 4(2):7-15.
Case Study 6
Derek Fawcett, Sridevi Brundavanam, Gérrard Eddy Jai Poinern. Growth
and corrosion Behaviour of amorphous micrometre scale calcium
phosphate coatings on Magnesium substrates. International Journal of
Materials Engineering. 2015; 5(1): 10-16.
6.3. Chapter Summary
58
Chapter 7 - General Discussion, Conclusions, and Further Work 59
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Acknowledgements
First and foremost I would like to express my sincere gratitude and thanks to my
principal supervisor Dr. Gérrard Eddy Jai Poinern, who possesses a great source of
knowledge. I am extremely honoured to have him as my supervisor and mentor. I would
like to thank him for his constant support and encouragement in my academic pursuits
and playing a huge role in making it all possible.
I am extremely obliged and thankful to Dr. Derek Fawcett for sharing his wealth of
knowledge and expertise during this project. His constant support and guidance during
this research work has been very valuable to me.
My sincere thanks to Dr. Xuan Thi Le for her help with FE-SEM in this study. I would
like to express my gratitude to the members of our research group, (Murdoch Applied
Nanotechnology Research Group) for sharing their knowledge, time and good humour
throughout the project work and making this period unforgettable in my university
career.
I would like to thank Associate Prof. Gamini Senanayake for his assistance in
performing the Electro chemical testing. I would like to thank Mr. Kenneth Seymour
providing training and assistance with XRD. Thanks are due to Mr. Peter Fallon for his
assistance with TEM.
I would like to express gratitude to Dr. Marc Hampton for providing training and
assistance with SEM. I would like to thank Mr. Andrew Foreman for providing training
and assistance with AAS.
I am extremely delighted to express my sincere gratitude and love to my parents
(Muralidhar & Bhargavi), parent-in-laws (Sampath Kumar & Rajya Lakshmi), my
brothers and my teachers for their constant encouragement and support throughout my
life’s journey.
Last but not least, I would like to acknowledge Murdoch University for providing a
scholarship and all other facilities for me to undertake my research studies in
nanotechnology.
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Chapter 1 - Introduction
1.1. Overview
The quality, health and wellbeing of human life are influenced by a number of factors
such as life style and medical fitness. However, the majority of the population, during
their life will experience medical problems ranging from damage, trauma and diseases,
and ultimately complications arising from aging tissues and organs. In most cases these
problems will often require surgical procedures to repair or replace the damaged or
diseased tissues or organs using available replacement materials or devices. In the case
of bone tissue, millions of people each year worldwide will require a bone transplant
and the source of the bone tissue comes from the patient or from a suitable donor. The
preferred source of bone tissue is from the patient and is known as an autograft. The
advantage of an autograft is its excellent biocompatibility, osteogenic properties and
furthermore delivers the patient’s own bone forming cells to the implant site. The
autograft technique has been used for a long time, delivers good clinical outcomes and
is considered the gold standard for bone transplantation clinical procedures.
Unfortunately, the source of potential donor sites is limited and a further complication is
donor site morbidity that has resulted in a search for alternative sources of bone tissue
[1-4]. Alternative materials such as allogenic bone grafts (sourced from another donor)
and xenograft bone grafts (sourced from another species) will result in a significant
response from the body’s immune system. A further complication associated with
allografts and xenografts is the serious threat of disease transmission [5]. Moreover,
ethical, medical and legal concerns involved in sourcing and using allografts and
xenografts has made their use difficult in many parts of the world. Importantly, all
2
autograft, allograft and xenograft procedures suffer from a limited supply of high
quality bone tissue that is suitable for medical use.
To address the shortcomings of conventional grafts, research and development efforts
have focused on using natural or creating biologically compatible synthetic materials
capable of repairing or replacing damaged or diseased bone tissues. A material or a
combination of materials that are used for this purpose are called biomaterials and can
be used to fabricate tissue scaffolds, implantable devices and prostheses. Biomaterial
development is a multidisciplinary field that covers a wide range of areas such as
material science, engineering, biotechnology, and biomedical sciences. And in recent
years there has been considerable effort made to produce a variety of engineered
implantable biomaterials capable of promoting the repair and regeneration of human
tissues such as bone, cartilage and skin [6-8]. Historically, metallic biomaterials have
been used for decades to replace damaged or diseased bone tissues in load-bearing
applications such as hip and knee replacements [9, 10]. Traditionally, metals used in
load-bearing applications have included stainless steels, cobalt chromium and titanium
alloys [11]. While materials such as polymers and ceramics with lower mechanical
strength and fracture toughness have been used in a variety of low load-bearing
applications [12-16]. In particular, because of their biocompatibility, calcium phosphate
ceramics have been used in a variety of applications such as bone repair, bone
augmentation, coating of metal implants and as filler material for both bone and teeth
[17-19]. Currently, metallic implants are extensively used in load-bearing applications
involving bone tissues.
3
Despite the many advantages offered by conventional metallic implants a number of
issues relating to unfavourable inflammatory responses and mechanical compliance still
exist. The adverse inflammatory response results from the slow release of toxic metallic
ions produced by corrosion and wear of the metallic implant into surrounding cells and
tissues [20, 21]. The ultimate detrimental result of metallic ion release is the significant
reduction of the implants biocompatibility [22]. The poor mechanical compliance
results from the significant difference in mechanical properties between metal implants
and surrounding bone tissue. The difference in these properties results in a clinical
phenomenon known as stress shielding that results in most of the load being carried by
the implant. The net result of stress shielding is the significant reduction in the load
being carried by surrounding bone tissues. The reduced load related stresses induce
bone resorption that ultimately leads to mechanical instability and failure of the implant
[23]. Another disadvantage of metal implants occurs when they are used as temporary
support structures such as pins, plates and screws during the repair of damaged bone. In
this case, a second surgical procedure is often needed to remove the implant after the
healing process has taken place. The second surgical procedure significantly increases
health costs and morbidity. These clinical issues highlight the need for new biologically
compatible materials that are not only capable of providing short-term mechanical
compliance and structural support during the healing process, but are also capable of
safely degrading with time and avoid the issues normally associated with conventional
metal implants.
Magnesium (Mg) is a lightweight, silvery-white metal and is the main constituent in a
number of Mg based alloys that are currently being used in aerospace and automotive
industries [24]. Because Mg’s high strength to weight ratio and mechanical properties
4
are similar to those of bone tissue, it is currently being investigated as a potential
alternative to conventional metals used in biomedical implants. For example, the density
of Mg is 1.74 g/cm3
while bone varies from 1.8 to 2.1 g/cm3 and the elastic modulus of
Mg (45 GPa) is within the modulus range of bone (40 to 57 GPa) [25, 26]. The close
similarity between the respective densities and elastic moduli have made Mg a potential
biomedical material capable of significantly reducing the effects of stress shielding and
bone resorption normally encountered in hard tissue engineering applications. Other
attractive features include biologically degradability and absorbability. In particular, the
release of Mg ions during corrosion could be considered beneficial since the body uses
Mg in a number of metabolic processes and apatite formation in bone [27]. The
biological importance of Mg is reflected by the fact that 30 g is stored in muscle and
bone tissues in an average adult body [28]. However, the major limitation that prevents
Mg’s widespread medical application as an implant material is its corrosion behaviour.
In the physiological environment, Mg’s corrosion rate is too fast to be an effective
biodegradable implant material. In addition, the body’s metabolic pathways cannot
safely handle the rapid formation of hydrogen gas in the physiological environment.
There are generally two possible routes to improve the corrosion resistance of Mg. The
first involves modifying the composition and microstructure and the second involves
producing a suitable surface treatment that protects the underlying metal. Most of the
compositional additives and surface treatments described in the literature were not
developed for medical applications and many of the additives and surface treatments are
capable of also producing toxic side effects [29]. Currently, there are no specifically
designed Mg based products commercially available in the biomedical sector for hard
tissue engineering.
5
The objective of this thesis was to develop biocompatible and biodegradable coatings
for Mg substrates capable of reducing and controlling the degradation rate and
increasing substrate biocompatibility for potential hard tissue engineering applications.
In recent years, the literature has shown a remarkable increase in the number of
publications investigating the development of biologically compatible coatings. For
tissue engineering applications, coatings should not only provide corrosion protection,
but they should also enhance bioactivity, biocompatibility and promote
osseointegration. And ideally, the coatings should also have antibiotic and drug delivery
properties. Importantly, the coatings must be able to biologically degrade at a controlled
rate and allow the underlying Mg substrate to slowly dissolve. Therefore, the coating
should only provide an effective protective barrier for a limited timeframe, thus,
allowing regenerating bone tissues to progressively replace the implant.
1.2. Scope of thesis
This thesis focuses on developing biodegradable and biocompatible calcium phosphate
coatings capable of extending the operational life of an Mg based implant and allow
sufficient time for tissue regeneration to take place. Calcium phosphates such as
hydroxyapatite, a mineral phase found in bone, are known to promote bone cell
adhesion, differentiation and osteointegration [30]. The main component of the
experimental work carried out in this thesis examines the types of calcium phosphates
formed on Mg substrates via chemical immersion and electrochemical techniques.
Advanced characterisation techniques such as X-ray diffraction (XRD), field emission
scanning electron microscopy (FESEM) and Fourier transform infrared spectroscopy
(FT-IR) were used to investigate the size, morphology, composition and topographical
features of the uncoated and coated substrates. Extensive corrosion studies were also
6
carried out on various coated substrate types in phosphate buffer saline (PBS) solution
and Ringer’s solution at 37 ºC to simulate body fluid conditions. Also investigated was
the adsorption capability of a representative coating, since a major function of bone is to
store minerals routinely used by the body.
1.3. Aims of thesis
The thesis is structured around four aims, each composed of one or two individual case
studies that allow a more detailed investigation into the various aspects of the research.
Aim 1. Chemically synthesize sufficient quantities hydroxyapatite from novel sources
for potential bone cements and coatings. The synthesized hydroxyapatite will have
similar chemical adsorption capabilities to natural bone. The synthesized hydroxyapatite
will also have similar bioactivity and biocompatibility properties to the mineral phase
found in bone (Case Study 1).
Aim 2: Investigate the chemical adsorption capabilities of hydroxyapatite (a
representative coating material) to determine its suitability as an effective bone
substitute during a tissue regeneration or tissue engineering procedure (Case Studies 2
& 3). This is of particular importance since the coating must be able to function like the
mineral phase found in bone.
Aim 3: Develop and use chemical immersion and electrochemical techniques to
synthesize calcium phosphate coatings. The use of advanced characterization techniques
such as X-ray diffraction (XRD) spectroscopy, scanning electron microscopy (SEM),
Energy Dispersive Spectroscopy (EDS) and Fourier Transform Infrared spectroscopy
7
(FT-IR) will be used to determine size, morphology, composition and topographical
features of the various substrate coatings (Case Studies 4, 5 & 6).
Aim 4: Conduct bio-corrosion studies to investigate the degradation behaviour of the
various coating types in phosphate buffer saline (PBS) solution and Ringer’s solution at
37 ºC and pH 7.4 to simulate body fluid conditions (Case Study 4 & 6). The results of
the corrosion studies will provide a guideline for future implant development.
The following chapters of this thesis address the above-mentioned aims, and elucidate
the current state of research in this field and its relevance to developing biodegradable
implants for hard tissue engineering applications.
References
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Elbow Surg. 16 (2007) S282-S285.
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transfer procedure, Arthroscopy, 20 (2004) e69-e73.
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6. A. Khademhosseini, J.P. Vancanti, R. Langer, Progress in tissue engineering, Sci
Am. 300 (2009) 64-71.
7. E.S. Place, N.D. Evans, M.M. Stevens, Complexity in biomaterials for tissue
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8. G. Poinern, R. Shackleton, S.I. Mamun, D. Fawcett, Significance of novel
bioinorganic anodic aluminum oxide nanoscaffolds for promoting cellular
response, Nanotechnol. Sci. Appl. 4 (2011) 11–24.
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10. P.T. Lee, D.L. Lakstein, B. Lozano, O. Safir, J. Backstein, A.E. Gross, Mid-to
long term results of revision total hip replacement in patients aged 50 years or
younger, Bone. Joint. J. 96-B (2014) 1047-1051.
11. M. Geetha, A.K. Singh, R. Asokamani, A.K. Gogia, TI based biomaterials, the
ultimate choice for orthopaedic implants – A review, Mater. Sci. 54 (2009) 397-
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12. Z.M. Huang, Y.Z. Zhang, M. Kotaki, S. Ramakrishna, A review of polymer
nanofibres by electrospinning and their applications in nanocomposites, Comp.
Sci. Tech. 63 (2003) 2223-2253.
13. J.C. Middleton, A.J. Tipton, Synthetic biodegradable polymers as orthopedic
devices, Biomaterials 21 ( 2000) 2335-2346.
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15. J.R. Liebermann, G.E. Friedlaender, ed. Bone Regeneration and Repair: Biology
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9
16. C.S. Cutter, B.J. Mehrara, Bone grafts and substitutes, J. Long Term Eff. Med.
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19. A. Stoch, W. Jastrze, E. Długoń, W. Lejda, B. Trybalska, G. Stoch, A.
Adamczyk, Sol-gel derived hydroxapatite coatings on titanium and its alloy
Ti6Al4V, J Mol Struct. 744 (2005) 633-640.
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Orthopaedic Implants, J. Bone Joint Surg. Am. 83 (2001) 428-428.
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Variation in cytokines induced by patients from different prosthetic materials,
Clin. Orthop. Relat. Res. 352 (1998) 323-230.
22. C. Lhotka, T. Szekeres, I. Steffan, K. Zhuber, K. Zweymüller, Four-year study
of cobalt and chromium blood levels in patients managed with two different
metal-on-metal total hip replacements, J. Orthop. Res. 21 (2003) 189-195.
23. J. Nagels, M. Stokdijk, P.M. Rozing, Stress shielding and bone resorportion in
shoulder arthroplasty, J. Shoulder. Elbow. Sur. 12 (2003) 35-39.
24. I.J. Polmear, Magnesium alloys and applications, Mater. Sci. Tech. 10 (1994) 1-
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25. A. Feng, Y. Han, The microstructure, mechanical and corrosion properties of
calcium phosphate reinforced ZK60A magnesium alloy composites, J. Alloys
Compd. 504 (2010) 585-593.
26. M. Razavi, M. Fathi, M. Meratian, Microstructure, mechanical properties and
biocorrosion evaluation of biodegradable AZ91-FA nanocomposites for
biomedical applications, Mater. Sci. Eng. A, 527 (2010) 6938-6944.
27. S.R. Kim, J.H. Lee, Y.T. Kim, D.H. Riu, S.J. Jung, Y.J. Lee, S.C. Chung, Y.H.
Kim, Synthesis of Si, Mg substituted hydroxyapatites and their sintering
behaviours, Biomaterials, 24 (2003) 1389-1398.
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Review of Material Properties, Surface Modifications and Potential as a
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11
Chapter 2 - Literature review
2.1. Overview and author contributions
Chapter 2 is composed of two peer-reviewed articles that survey the current literature
and examine the biomedical potential of using magnesium (Mg) substrates for hard
tissue bioengineering applications. The first review paper provides a brief history of
Mg, its physical, chemical and mechanical properties, and its use in both medical and
industrial applications. In particular, its high strength to weight ratio has made it an
attractive material in industries such as transport and aerospace. The review also
examines Mg biological performance and its potential application as biodegradable
orthopaedic implants. However, Mg’s high corrosion rate in the physiological
environment is its main disadvantage and currently prevents its use in many
bioengineering applications. The metal’s low corrosion resistance, especially in
electrolytic aqueous environments such as the physiological environment, where it
rapidly degrades must be addressed. In recent years there has been a significant effort to
improve Mg’s corrosion resistance and to slow down its degradation rate in a variety of
environments. In particular, the review focuses on biological corrosion and discusses
the various factors influencing Mg’s corrosion rate such as alloying elements, surface
modifications and surface treatments.
The second peer-reviewed article focuses on using chemical immersion techniques for
producing protective calcium phosphate coatings on Mg substrates. Chemical
immersion is an economic, efficient and straightforward technique that offers a direct
method of depositing calcium phosphate coatings on Mg substrates. Besides reducing
the corrosive effects of the physiological environment, the coatings also have the
12
potential to significantly improve biocompatibility and promote bone formation at the
coating-substrate interface of the Mg based implant.
The author contributions consisted of G.E.J. Poinern acting as principal supervisor who
designed the overall concept for each review paper with S. Brundavanam being the
major contributor to the papers. In the second review paper S. Brundavanam also acted
as first author and significantly contributed to the content of the paper. All text, tables
and images were carried out by S. Brundavanam under the supervision of D. Fawcett. S.
Brundavanam was assisted by G.E.J. Poinern and D. Fawcett in over-coming some of
the various technical difficulties encountered during the paper preparation and with the
editorial changes to the manuscript as recommended by reviewers. All authors provided
feed-back during the preparation of the paper which was coordinated by S.
Brundavanam.
2.2. Published Review Articles
Review Article 1
Gérrard Eddy Jai Poinern, Sridevi Brundavanam and Derek Fawcett, Biomedical
Magnesium Alloys: A Review of Material Properties, Surface Modifications and its
Potential as a Biodegradable Orthopaedic Implant. American Journal of
Biomedical Engineering. 2012; 2 (6); 218-240.
Review Article 2
Sridevi Brundavanam, Gérrard Eddy Jai Poinern, Derek Fawcett, Chemical
immersion coatings to improve biological degradability of magnesium substrates
for potential orthopaedic applications. International Journal of Biomedical
Materials Research. 2014; 2(2): 7-14.
American Journal of Biomedical Engineering 2012, 2(6): 218-240 DOI: 10.5923/j.ajbe.20120206.02
Biomedical Magnesium Alloys: A Review of Material Properties, Surface Modifications and Potential as a
Biodegradable Orthopaedic Implant
Gérrard Eddy Jai Poinern*, Sridevi Brundavanam, Derek Fawcett
Murdoch Applied Nanotechnology Research Group, Department of Physics, Energy Studies and Nanotechnology, School of Engineering and Energy, Murdoch University, Murdoch, Western Australia, 6150, Australia
Abstract Magnesium and magnesium based alloys are lightweight metallic materials that are extremely biocompatib le and have similar mechanical properties to natural bone. These materials have the potential to function as an osteoconductive and biodegradable substitute in load bearing applicat ions in the field of hard t issue engineering. However, the effects of corrosion and degradation in the physiological environment of the body has prevented their wide spread applicat ion to date. The aim of this review is to examine the properties, chemical stability, degradation in situ and methods of improving the corrosion resistance of magnesium and its alloys for potential application in the orthopaedic field. To be an effective implant, the surface and sub-surface properties of the material needs to be carefully selected so that the degradation kinetics of the implant can be efficiently controlled. Several surface modification techniques are presented and their effectiveness in reducing the corrosion rate and methods of controlling the degradation period are discussed. Ideally, balancing the gradual loss of material and mechanical strength during degradation, with the increasing strength and stability of the newly forming bone tissue is the ultimate goal. If this goal can be achieved, then orthopaedic implants manufactured from magnesium based alloys have the potential to deliver successful clinical outcomes without the need for revision surgery.
Keywords Magnesium, Bio logical Corrosion, Biocompatibility, Alloys, Surface Modification
1. Introduction The skeletal system of the human body is a complex three-dimensional structure that is important for two main reasons. The first arises from the need to structurally support the many body organs and other related tissues. The second is the attachment of the numerous muscle groups that are needed for body movement and locomotion. The skeleton is constructed of two types of tissue, the first is a hard t issue called bone and the second is a softer tissue composed of cartilag inous materials. The adult human skeleton consists of 206 bones[1]; some provide protection to the internal organs, while others perform specialized functions such as transmitting sound vibrations in the inner ear. The bone matrix also provides a natural reservoir for cells and mineral ions that play an important role in maintain ing the biochemical balance within the body. For example, calcium is an important element involved in muscular action and nerve conduction and its level in the body is closely monitored and regulated by a process called homeostasis[2].
*Corresponding author: [email protected] (Gérrard Eddy Jai Poinern) Published online at http://journal.sapub.org/ajbe Copyright © 2012 Scientific & Academic Publishing. All Rights Reserved
Bone is a natural two phase organic-inorganic ceramic composite consisting of collagen fibrils with an embedded inorganic nano-crystalline component. The primary organic phase of the bone matrix is Type I collagen, which is secreted by osteoblast cells to form self-assembled fibrils[3, 4]. The fibrils are bundled together and orientate themselves parallel to the load-bearing axis of the bone. The fibrils are typically 300 nm long, develop a 67 nm periodic pattern in which a 40 nm gap or hole is formed between the ends of the fibrils and the remain ing 27 nm overlaps the bundle behind[5]. This pattern creates discrete and discontinuous sites for the deposition of plate-like nanometre sized hydroxyapatite (HAP) crystals, which forms the second phase of the bone matrix. HAP is a mineral predominantly composed of calcium phosphate which has the general chemical formula of[Ca10 (OH) 2(PO4)6]. It is the main inorganic component of bone and teeth, accounting for up to 65% by weight of cortical bone and in the case of teeth it accounts for 97 % by weight of dental enamel in mammalian hard tissue[6]. The discontinuous discrete sites limit the growth of the HAP crystals and force the crystals to grow with a specific crystalline orientation which is parallel to the load-bearing axis of the bone and collagen fibrils. The crystal p lates typically have a length of 50 nm, a width of around 25 nm and on average a thickness of 3 nm[7-10]. The HAP also has
219 American Journal of Biomedical Engineering 2012, 2(6): 218-240
trace amounts of potassium, manganese, sodium, ch loride, hydrogen phosphate, citrate and carbonate[11]. The final component of the bone matrix consists of the non-collagen organic proteins such as the phosphor-protein group which are believed to regulate the formation of the inorganic crystal phase by influencing the size, orientation and the depositional environment within the spaces between the collagen fibrils. The phosphor-protein group is also believed to be the source of calcium and phosphate ions used in the formation of the mineral phase[12].
The organic phase gives bone its flexib ility, while the inorganic phase provides bone with its structural rig idity[13, 14]. The incorporation of organic and inorganic phases in the matrix g ives bone its unique mechanical properties such as toughness, strength, and stiffness. It is the combination of these properties that give bone and the skeletal system in general, its remarkable ability to withstand the various mechanical and structural loads encountered during normal and intense physical activity[15]. However, not all bone tissue in the body has the same properties and this is characterized by the presence of two types of bone. The first type consists of a hard outer layer of compact (cortical) t issue, while the second type forms the less dense and spongy (trabecular) tissue which fills the interior of the bone. This spongy interior contains marrow and the many blood vessels that supply nutrients and remove waste products from the bone tissues. Both the cortical bone and the trabecular bone are composed of the same organic and inorganic phases discussed above, but they differ in the amount of each phase present. The two bone types also differ in their respective porosities and in their structural arrangement. The amount of cortical and trabecular tissue found in bone is dependent on the external load being applied and the frequency of the load[16]. Despite its remarkable mechanical and structural properties bone can fracture from three main causes: 1) a fracture caused by sudden injury; 2) Fatigue or stress fractures resulting from repeated cyclic loads; and 3) Pathological fractures resulting from bone infections and tumours[17]. The surgical implantation of artificial biomaterials of specific size and shape is an effective solution in restoring the load bearing capacity and functionality of damaged bone tissue. The design and selection of biomaterials is highly dependent on the specific medical applicat ion. Therefore, it is imperative that new biomaterials being developed for load bearing orthopaedic implant applications should have excellent biocompatibility, comparable strength to natural bone, and produce no cytotoxicity effects[18, 19].
Metallic b iomaterials have been used since the early 1900s to replace damaged or diseased hard tissues. And as early as 1907, a magnesium alloy was used by Lambotte, to secure a bone fracture in the lower leg[20, 21]. Metallic implants are generally used in load bearing applications where their high mechanical strength and fracture toughness make’s them superior to ceramics, polymeric materials and polymer / ceramic composites. Metallic implant materials currently used include stainless steel, cobalt-chrome alloys and
titanium and its alloys. At present there are two major problems associated with using the metallic implants. The first involves the mis match between the mechanical properties of the metallic alloy and the surrounding natural bone tissue. The elastic modulus of both stainless steel and cobalt-chrome alloys is around ten times greater than that of bone, while a titanium alloy such as Ti-6Al-4V is around five times greater[22]. Bone tissue is constantly undergoing remodelling and modification in response to imposed stresses produced by normal everyday activities. The mechanical mis match between bone and different metallic implant materials results in a clin ical phenomenon known as stress shielding. The stress-shielding phenomenon occurs when the implant carries the bulk of the load and the surrounding bone tissue experiences a reduced loading stress. The reduced loading stress experience by the surrounding bone tissue ultimately leads to bone resorption[23, 24]. The second problem stems from mechanical wear and corrosion of the implant and results in the release of toxic metallic ions such as chromium, cobalt and nickel into the body. These harmful metallic ions solicit an inflammatory response from the body’s immune system and the surrounding tissues which reduces the biocompatibility of the implant[25, 26, and 27]. This is in total contrast to the corrosion products of magnesium (Mg) which can be considered physiologically beneficial, with the adult body storing around 30 g of Mg in both muscle and bone tissue[28]. The importance of Mg to the body stems from the fact it is bivalent ion which is used to form apatite in the bone matrix and is also used in a number of metabolic processes within the body[29]. And recently, Robinson et al. reported the novel antibacterial properties of Mg metal against Escherichia coli, Pseudomonas aeruginosa and Staphylococcus aureus[30].
Mg is a lightweight, silvery-white metal that is relatively weak in its pure state and is generally used as an alloy in engineering applications. The density of Mg and its alloys are around 1.74 g/cm3 at 20oC, which is 1.6 and 4.5 times less dense than alumin ium and steel, respectively[31]. Interestingly, the density of Mg is slightly less than natural bone which ranges from 1.8 to 2.1 g/cm3, while the elastic modulus of pure Mg is 45 GPa and human bone varies between 40 and 57 GPa[32, 33 and 34]. Because of this close similarity in the respective elastic moduli, using Mg in hard tissue engineering applications would greatly reduce the possibility of stress shielding and prevent bone resorption. Thus, Mg with its similar mechanical properties to natural bone, combined with its biocompatibility, makes it a promising material for the development of biodegradable orthopaedic implants[33, 35].
Polymeric materials have also been used in a number of tissue engineering applications since they have many attractive properties such as being lightweight, ductile in nature, biocompatible and biodegradable. Polymers are materials with large molecules composed of small repeating structural units called monomers. The monomers are usually attached by covalent chemical bonds, with cross-linking taking place along the length of the molecule. It is the
Gérrard Eddy Jai Poinern et al. : Biomedical Magnesium Alloys: A Review of Material Properties, 220 Surface Modifications and Potential as a Biodegradable Orthopaedic Implant
amount of cross-linking that gives the polymer its physiochemical propert ies. Many polymeric materials have been investigated since the body’s natural processes can easily handle the by-products resulting from their degradation, with the by-products being easily excreted in the urine. Natural polymers such as polysaccharides[36-40], chitosan[41-46], hyaluronic based derivatives[47-50] and protein based materials such as fibrin gel[51, 52] and collagen[53-56], have all produced favourable outcomes in a number of tissue engineering applications.
Similar studies using synthetic biopolymers composed of simple high purity constituent monomers, fabricated under controllable formation conditions have produced a variety of tissue scaffolds and implants with tuneable and predictable physio-mechanical properties. These biopolymers also have low toxicity reactions with the body and their degradation rate can be easily controlled. Examples of synthetic biodegradable polymers include Poly (lact ic acid), PLA[57-62], Po ly (L-lactic acid), PLLA [63-66], Po ly (lactic-co-glycolic acid), PLGA[67-70], Poly-capro lactone PCL[71-74] and Poly (g lycolic acid ), PGA [75-78]. These biopolymers are generally poly-α-hydrox esters that de-esterifies in the body as the polymer degrades to simple metabolites[79]. Currently available b iodegradable sutures in clinical use are made from PLA and PGA. These synthetic biopolymers can also be made into different shapes and structures, such as pellets, rods, disks, films, and fibres as required for the specific applicat ion. Some of these applications include biodegradable sutures, bone and dental cements, bone grafting materials, plates, screws, pins, fixation devices and low load bearing applications in orthopaedics[80, 81]. However, even with their many attractive properties, biopolymers have low mechanical strength when compared to ceramics and metals, which has resulted in them being used in soft tissue reconstruction and low-load bearing applications. The major advantage that Mg and its alloys have over biopolymers is its superior mechanical strength, which is typically double that of biopolymers.
Ceramics are non-metallic, inorganic materials that are used in hard tissue engineering applications where they are collectively termed bioceramics. The important properties of bioceramics that make them highly desirable fo r b iomedical applications are: 1) they are physically strong; 2) they are both chemically and thermally stable; 3) they exh ibit good wear resistance, and 4) they are durable in the body environment[82]. In addit ion, they are readily available, can be shaped to suit the application, they are biocompatible, hemocompatible, nontoxic, non-immunogenic and can be easily sterilised[83]. But unlike Mg and its alloys, bioceramics such as HAP, tend to be brittle, have low fracture toughness and are not as resilient. However, some bioceramics have found application in hip jo ints, coatings on implants, maxillofacial reconstruction, bone tissue engineering and drug delivery devices[81, 84-86].
A composite material consists of two or more distinct parts or phases[85]. The major advantage of using a composite biomaterial stems from the fact a single-phase material may not have all the required properties for a particular applicat ion[86]. However, by combin ing one or more phases with differing physical and chemical properties it is possible to create a composite material with superior properties to those of the indiv idual components. A good example of a natural composite is bone, which is a composed of Type 1 collagen and HAP. A typical manmade example of a biomedical composite is a b ioactive coating of HAP or a bioactive glass deposited on to the surface of a titanium implant to promote bone attachment[87]. Composites, such as a 2-phase HAP-polymer mixture have also been developed to create a biomaterial with similar p roperties to natural bone for hard tissue engineering applications[88]. Unfortunately, as mentioned above, biopolymers biodegrade with time and as a result, the load bearing capacity and fracture toughness of the implant will decline with time.
When comparing the propert ies of Mg and its alloys with metals, polymers, ceramics and composites it can be shown that Mg and its alloys have many properties that are comparable, if not superior, see Table 1. However, despite its many advantages, Mg has the disadvantage of having a high corrosion rate in the body. And as a result, medical application of Mg based implants has been severely limited due to the electrolytic aqueous environment of the chloride rich body fluid (pH ranges between 7.4 and 7.6). Furthermore, there are two serious consequences of the rapid corrosion rate of Mg implants. The first is the rap id evolution of subcutaneous hydrogen gas bubbles which are produced at a rate too high fo r the surrounding tissues to handle[89, 90]. These bubbles usually appear within the first week after surgery and can be easily treated by drawing off the gas using a subcutaneous needle[91]. The second consequence of the high corrosion rate is the loss of mechanical integrity of the Mg implant being used in the load bearing application. The rapid decrease in mechanical properties resulting from exposure to the body fluid environment means that the implant is unable to provide the necessary support for the healing bone tissue. Generally, the implant would be expected to maintain its mechanical integrity between 12 to 18 weeks while the healing process takes place and then slowly degrade while natural bone tissues replace the implant[92].
This article reviews the biological performance, mechanical properties and potential applicat ion of biodegradable Mg based alloys for orthopaedic implants. The major disadvantage of using Mg in many engineering applications is its low corrosion resistance, especially in electrolytic, aqueous environments where it rapidly degrades. To slow the degradation rate in situ, factors influencing the corrosion rate such as alloying elements, surface modification and surface treatments are examined and discussed in the following sections.
221 American Journal of Biomedical Engineering 2012, 2(6): 218-240
Table 1. Some mechanical properties of selected materials
Tissue/Material Density
(g cm-3)
Compressive Strength (MPa)
Tensile Strength (MPa)
Elastic Modulus
(GPa)
Natural Materials Arterial wall Collagen Collagen(Rat tail tendon) Cancellous bone Cortical bone
- - -
1.0 – 1.4 1.8 – 2.0
- - -
1.5 – 9.3 160 Trans. 240 Long.
0.50 -1.72
60 -
1.5 – 38 35 Trans. 283 Long.
0.001 1.0
3.75 – 11.5 0.01 – 1.57
5 - 23
Magnesium Alloys Pure magnesium AZ31 (Extruded) AZ91D (Die cast)
1.74 1.78 1.81
20 - 115 83 - 97
160
90 - 190
241 - 260 230
45 45 45
Other metal alloys Cobalt-Chrome Alloys Stainless Steel Titanium Alloys
7.8 7.9 4.4
- - -
450 - 960 480 – 620 550 – 985
195 - 230 193 – 200 100 – 125
Ceramics Synthetic- Hydroxyapatite Alumina Ceramics (Al2O3 80% - 99%)
3.05 – 3.15
3.30 – 3.99
100 - 900
2000 – 4000
40 – 200
-
70 – 120
260 – 410
Polymers Polymethylmethacrylate (PMMA) Polyethylene- terephthalate (PET)
1.12 – 1.20
1.31 – 1.38
45 – 107
65 – 90
38 – 80
42 – 80
1.8 – 3.3
2.2 – 3.5
Note: Table compiled from references[122, 126, 213, 218, 219, 220 and 221]
2. Biological Corrosion of Magnesium 2.1. Corrosion Mechanism
When unprotected chemically pure magnesium is exposed to humid atmospheric air it develops a thick dull gray amorphous layer composed of magnesium hydroxide[Mg (OH)2]. The oxidation rate of this protective oxide layer is typically around 0.01 mm/yr, while the oxidation rate in salt water is around 0.30 mm/yr[93]. In magnesium alloys, controlling the alloying chemistry and the overall microstructure of the alloy can significantly reduce the corrosion rate.
Table 2. Corrosion rates for some magnesium alloys immersion in various media
Material
In vitro corrosion rate (mg.cm-2.h-1)
In vivo corrosion rate
(mg.mm-2.yr-1) Hanks Solution
Simulated Body Fluid
Pure Mg (99.95%) 0.011 0.038 -
AZ31 0.0065 - 1.17 AZ91 0.0028 - 1.38
LAE442 - - 0.39 WE43 - 0.085 1.56
Note: Table compiled from references[92, 96, 126, 213, 214 and 215]
For orthopaedic applications pure magnesium finds the
human body a highly aggressive corrosive environment, see Table 2. The body flu ids are composed of water, dissolved oxygen, proteins and electro lytic ions such as chloride and hydroxide. In this environment, magnesium with a negative electrochemical potential o f -2.37 V, is very susceptible to corrosion and results in free ions migrat ing from the metal surface into the surrounding fluid environment.
These ions can form chemical species, such as metal oxides, hydroxides, chlorides and other compounds. In thermodynamic terms, with the assumption that there is no barrier to oxidation of the metal surface, the reaction would be very rap id, evolv ing hydrogen gas and consuming the metal substrate surface. But in reality the electrochemical reaction results in the migration of ions from the metal surface into solution, which forms species that result in the formation of an oxide layer that adheres to the metal surface. The Mg (OH)2 layer fo rmed on the metal surface is slightly soluble and reacts with chorine ions to form highly soluble magnesium chloride and hydrogen gas[94, 95]. When the oxide layer fully covers and seals the metal surface, it forms a kinetic barrier or passive layer that physically limits or prevents further migration of ionic species across the metal oxide solution interface.
The corrosion of Mg in an aqueous physiological environment can be expressed in the following equations. The primary anodic reaction is expressed by the partial reaction presented in equation (1), at the same time the reduction of protons is expressed by the partial reaction occurring at the cathode (2).
Gérrard Eddy Jai Poinern et al. : Biomedical Magnesium Alloys: A Review of Material Properties, 222 Surface Modifications and Potential as a Biodegradable Orthopaedic Implant
Anodic reaction: Mg → Mg2- + 2e- (1) Cathodic reaction: 2H2O + 2e- → 2OH- + H2 (2)
Another undesirable consequence of the corrosion process in Mg and its alloys is the formation of hydrogen gas. The rapid formation o f hydrogen gas resulting from the rich chorine environment produces subcutaneous gas bubbles, which generally appear within the first week after surgery and then disappear after 2 to 3 weeks[92]. During the init ial gas formation a subcutaneous needle can be used to draw off the gas. In 2007, Song postulated that a hydrogen evolution rate of 0.01 ml/cm2/day can be tolerated by the human body and does not constitute a serious threat[96].
If the Mg corrosion rate can be regulated so that the hydrogen evolution rate is below this value, then the implant will not create a gas threat. The reactions of solid Mg and the Mg (OH)2 layer with chorine ions in the aqueous environment are presented in equations (3) and (4).
Solid Mg: Mg (s) + 2Cl-(aq) → MgCl2 + 2e- (3)
Mg (OH)2 layer: Mg (OH)2 (s) + 2CL-
(aq) → MgCl2 + 2OH- (4) The general reaction of the corrosion process is presented
in equation (5). Mg (s) + 2H2O (l) → Mg (OH)2 (s) + H2 (g) (5)
Corrosion in the aqueous environment of the body is not as straight forward as corrosion in the industrial environment. This is due to the corrosion rate being influenced by a variety of other factors such as: 1) the pH of body fluids; 2) variations in the pH value; 3) concentration of ions; 4) the presence of proteins and protein adsorption on the orthopaedic implant; and 5) the influence of the surrounding tissues[97, 98 and 99].
2.2. Types of Biological Corrosion
An important property of the oxide layer is its ability to remain fixed to the metal surface during a variety of mechanical loading situations. If the oxide layer ruptures during mechanical loading it will expose the pure Mg substrate to body fluids which will result in further corrosion. The clinical repercussion of the corrosion process is the loss of mechanical strength and the ultimate failure of the implant. Typical forms of Mg corrosion encountered within the body environment are d iscussed in the following sections.
2.2.1. Galvanic Corrosion
Galvanic corrosion takes place between two dissimilar metals, each with a different electrochemical potential, when they are in contact in the presence of an electrolyte which provides a pathway for the transfer of electrons. The less noble metal becomes anodic, corrodes and produces a build up of corrosion by-products around the contact site. For example, if gold screws are used to attach an Mg plate to bone during reconstructive procedure, the resulting electrolytic effect of the body fluids (serum or interstitial flu id) would preferentially attack the Mg plate; see Figure 1[100]. Therefore, it would be good design practice to use
metals with similar electrochemical properties when designing implant devices. For example, the fixation screws used to attach an Mg plate during a bone reconstruction procedure should be made of a titanium (Ti) alloy, since Ti is the closest metal to Mg in the electrochemical series. Mg is the most reactive metal in the electrochemical series and will always be the anode in any corrosion reaction[101]. Therefore, selection of Ti alloy fixation screws to secure the Mg plate ensures the lowest possible corrosion rate. Galvanic corrosion can also result from the presence of inter-metallic alloying elements or impurit ies present in the Mg matrix, see Figure 2.
Figure 1. Galvanic corrosion between dissimilar metals
Figure 2. Galvanic corrosion resulting from inter-metallic elements
2.2.2. Granular Corrosion
In many metal alloys, inter-granular corrosion can occur from the presence of impurities and inclusions which are deposited in the grain boundary regions during solidification. Following solidification, numerous galvanic reactions takes place between the metal matrix and the various impurities and inclusions. The ensuing corrosion rate at the various grain boundary regions exceeds that of the grains and results in an accelerated corrosion rate of the metal matrix. However, in the case of Mg alloys, inter-granular corrosion does not occur since the grains tend to be anodic, while their boundaries are cathodic in nature compared to the interior of the grains. The resulting grain boundary corrosion undercuts nearby grains which subsequently fall out of the matrix[102].
223 American Journal of Biomedical Engineering 2012, 2(6): 218-240
2.2.3. Pitt ing Corrosion
Pitting corrosion of Mg results from the rapid corrosion of small-localized areas which damage the protective surface oxide layer; see Figure 3. This form of corrosion is more serious than other forms of corrosion since the surface pits are difficu lt to see due to the presence of corrosion products. The pits are small, highly corrosive and continue to grow downwards, perforating the metal matrix[103]. After init ial nucleation at the surface, the presence of impurit ies in the Mg alloy microstructure often assists in further corrosion due to the galvanic differences in the materials[104, 105]. The environment within the pit is very aggressive, with chlorides species from the body flu ids and Mg+ ions from anodic dissolution greatly aggravating the situation. In addition, the mouth of the pit is s mall and prevents any dilution of the pit contents, which adds to the accelerating autocatalytic growth of the pit. During this process, electrons flowing from the pit make the surface surrounding the pit entrance become cathode-protected and the protective oxide layer is further weakened. Once pitting starts, an Mg component can be totally penetrated within a relatively short period of time and in the case of a biomedical implant, its load bearing capacity would be greatly reduced to the point of failure. Another problem associated with pitting arises from localised increase in stress produced by the pit, which has the potential to form cracks[106]. The fo rmation of stress corrosion cracking and metal fatigue cracks in the pits can lead to failure of the implant during normal loading conditions.
Figure 3. Pitt ing corrosion site at the surface of a magnesium component
Figure 4. Crevice corrosion occurring between magnesium components in a body fluid environment
2.2.4. Crevice Corrosion
Crev ice corrosion is local contact corrosion that occurs between metal and metal/non metal components. For example, if a magnesium plate is to be fixed in location by a set of screws with a small gap between the screw head and plate. The gap must have sufficient width to allow the flow of the body fluids through the gap and prevent any stagnant flow, see Figure 4. The stagnant flow results in the build up of Mg+ ions, with an Mg+ ion concentration gradient soon set up between the entrance and the dead end of gap. The subsequent corrosion cell then starts to attack the metal components of the implant[107].
2.2.5. Fretting Corrosion
Fretting corrosion is the result of damage p roduced by metal components in direct physical contact with each other in the presence of small vibratory surface motions. The micro-motions are produced by normal every day activities experienced by the human body which result in mechanical wear and metallic debris between the surfaces of metal components making up the biomedical implant[108]. During daily activ ity, the micro-motions remove the passive surface layer o f the metallic components in direct contact, exposing fresh metal underneath. Then both the fresh metal surfaces and the metallic surface debris undergo oxidation. The surface debris has a further detrimental effect by acting as an abrasive agent during subsequent micro-motions. The corrosion rate is dependent on the applied load, the resulting fretting motion, the microstructure of the metal o r metal alloys used in the implant and solution chemistry in the region around the fretting zone[109, 110]. During the corrosion process metallic ions are produced which can form a wide range of organic-metallic complexes and some metallic implants can release toxic metallic ions such as chromium, cobalt and nickel. These harmfu l metallic ions significantly reduce the biocompatibility of the implant and solicit a major inflammatory response from the body’s immune system[25, 26, and 27]. In the case of magnesium, metallic ions released during fretting, can be considered physiologically beneficial since these ions can be consumed or absorbed by the surrounding tissues, or be dissolved and readily excreted through the kidneys. Fretting corrosion is common in load bearing surfaces and is also capable initiat ing fatigue cracks in the fretting zone. Once formed the crack can propagate into the bulk of the metal matrix and can lead to the failure of the implant.
2.2.6. Erosion Corrosion
Erosion corrosion occurs from the wearing away of the metal surface or passive layer by the impact of wear debris in the body environment surrounding the implant. The metallic debris impacts on the surface of the implant, transferring energy into the region of the collision and plastically deforming the surface. During the deformation p rocess the surface becomes work harden to the point where the next impact exceeds the strain required fo r surface fracturing,
Gérrard Eddy Jai Poinern et al. : Biomedical Magnesium Alloys: A Review of Material Properties, 224 Surface Modifications and Potential as a Biodegradable Orthopaedic Implant
pitting or chip format ion. With the passage of time, the numerous impacts result in material loss from the metal surface[111]. For example, a femoral head of a Cobalt-Chromium implant will have numerous scratches after 17 years of implantation in a patient[112]. A ll bio-metals used in implants inevitably corrode at some fin ite rate when immersed in the complex electrolytic environment of the body; even Ti alloys with the lowest corrosion rate produce corrosion debris. The debris can significantly influence the wear behaviour and erosion resistant properties of the implant. However, the effects of erosion may not be noticed until there is a significant loss of metal which ultimately leads to the clinical failure of the implant.
2.2.7. Stress Corrosion
When an electrochemical potential is formed between stressed and unstressed regions of a metal implant under load, there is an increase in the chemical act ivity of the metal. This stress initiated corrosion mechanism effect ively increases the corrosion rate, usually by two to three times above the normal uniform rate. Th is usually results in the formation of s mall cracks that concentrate stress within the loaded implant, a mechanism know as stress corrosion cracking (SCC). Mg SCC can occur in any load stressed implant immersed in the dilute chloride environment of the body fluids. SCC init iated cracks grow rapidly and extend between the grains throughout the metal matrix[113, 114]. The progress of SCC is also influenced by the strain rate resulting from the implant loading cycles and the presence of hydrogen gas produced by the corrosion process[115, 116]. Current research suggests that chloride ions produce pitting in the protective surface layer, which ultimately leads to a break down in the layer exposing the underlining Mg matrix to the electrolytic flu ids of the body environment. The resulting hydrogen diffuses into the stressed zone of the metal matrix ahead of the crack tip and allows the SCC crack to advance through the zone[117-119]. Fracture and failu re of the implant will occur when the SCC is below the normal operating stress of the implant.
2.2.8. Corrosion Fatigue
Corrosion fatigue is the result of a material being exposed to the combined effects of a cycling load and a corrosive environment[120]. In general, metal fatigue is the damage caused by the repeated loading and unloading of a metal component. The cyclic stress initiates the formation of microscopic cracks on the metal surface and also damages the protective passive layer. If there are any surface imperfections such as pores or pitting from corrosion, they become crack nucleat ion sites which can significantly speed up crack g rowth rates. In the body’s environment the cracks become localized electrochemical cells that promote further corrosion. Mg in particular is susceptible to corrosion fatigue due to the presence of chloride ions in the body fluids. Corrosion within the crack promotes crack propagation and in combination with cyclic loading, the crack growth rate significantly increases. Eventually the loading stress exceeds
the SCC threshold and the crack grows to a critical size resulting in the fracture of the metallic implant. The body environment can significantly reduce the fatigue life of Mg alloys, producing lower failure stresses and considerably shorter failure times.
3. Magnesium and its Alloys For biomedical applicat ions, the composition of the
material being considered is a crucial factor since many of the elements that make up commercially available materials for industrial applications are extremely toxic to the human body. Therefore, in addition to meeting the mechanical properties needed for a particular biomedical applicat ion, the material must also be biocompatible. Ideally, a biodegradable biomedical device should be composed of materials or alloys that are non toxic or carcinogenic. It would also be very advantageous if the material was composed of elements and minerals already present and compatible within the body such as magnesium, calcium and zinc, see Tab le 3. Furthermore, the material should have a controllable d issolution rate or slow corrosion rate that permits the biomedical device or implant to maintain its mechanical integrity until the surrounding tissues heal and are capable of carrying the load once again. After the healing process has taken place, the load bearing properties of the biomedical implant are no longer required and the implant material should then be able to slowly dissolve away. Furthermore, the resultant by-products of the degradation process should be non-toxic; capable of being consumed or absorbed by the surrounding tissues, or being dissolved and readily excreted through the kidneys. Thus, for Mg and its alloys to be used as an effective biodegradable implant it is necessary to control their corrosion behaviour in the body flu id environment[121].
3.1. The Influence of Alloying Elements on Physical and Mechanical Properties
There are three major groups of Mg alloys: the first group consists of pure Mg; the second group consists of aluminium (Al) containing alloys such as AZ91, AZ31 and rare earth elements (RE) such as AE21; and the final group consists of the Al free alloys such as Mg-Ca, W E, MZ and WZ. The use of alloying elements such Al, Ca, Li, Mn, Y, Zn, Zr and RE in Mg alloys can significantly improve the physical and mechanical properties of the alloy by: 1) refin ing the grain structure; 2) improving the corrosion resistance; 3) form inter-metallic phases that can enhance the strength; and 4) assist in the manufacture and shaping of Mg alloys.
Impurit ies commonly found in Mg alloys are Be, Cu, Fe and Ni and the levels of theses impurities are restricted to within specific limits during the production of the alloy, see Table 3. The range of acceptable levels for Be ranges from 2 to 4 ppm by weight, while Cu is (100-300 ppm), Fe (30-50 ppm) and Ni (20-50 ppm)[122]. Since both Be and Ni are carcinogenic, their use in b iomedical applications should be
225 American Journal of Biomedical Engineering 2012, 2(6): 218-240
avoided as alloying elements. While elements such as Ca, Mn and Zn are essential trace elements for human life and RE elements exhib iting anti-carcinogenic properties should be the first choice for incorporation into an alloy. Studies by Song have suggested that very small quantities of RE elements and other alloying metals such as Zn and Manganese (Mn) could be tolerated in the human body and could also increase corrosion resistance[123]. Mn is added to many commercial alloys to improve corrosion resistance and reduce the harmful effects of impurit ies[124]. Mg alloys containing rare earth elements have also been found to increase the resistance to the flow of Mg2+ ions out of the Mg matrix via the Mg oxide layer[125]. During the degradation process the RE elements remained localised in the corrosion layer, which also contained high levels of both calcium and phosphorous. Also during this period a thin amorphous calcium phosphate layer formed over the surface of the oxide layer[92, 126].
Recent studies by Witte et al. have investigated the degradation behaviour of Mg based alloy rods and polymer based control rods[poly (lactic acid)] in animal models. Rods of 15 mm d iameter and 20 mm long were inserted into the femur of guinea p igs and the rods degradation profile monitored. The percentage compositions by weight of the Mg alloys investigated consisted of two alumin ium-zinc alloys composed of 3% Al and1% Zn AZ31 and 9% Al and 1% Zn AZ91 with the balance of the alloys composed of pure Mg. In addition two RE alloys were studied, the first consisted of 4% yttrium and a 3% rare earth mixture composed of neodymium, cerium and dysprosium W E43 and the second composed of 4% lithium, 4%, aluminium and a 2% rare earth mixture o f cerium, lanthanum, neodymium and praseodymium LAE442[92, 127]. The implants were harvested at 6 and 18 weeks, with complete implant degradation occurring at 18 weeks. During this time radiographs were regularly taken, while a micro -tomography-based technique using X-ray synchrotron
radiation was used to characterize the implant’s degradation process. All Mg based alloy implants were found to be beneficial and promoted new in situ bone tissue format ion, while the polymer control rods produced a less significant effect. The LAE442 alloy had the greatest resistance to corrosion, while the other alloys all had similar, but lower values of corrosion resistance and degraded at similar rates[92].
While Mg is potentially an ideal biocompatible implant material due to its non-toxicity to the human body, the safe long term use of an Mg based alloy needs to be carefully studied. Magnesium based alloys have also been used in vivo; for example an AZ91 alloy rods were implanted into the femur of a number of rabbit models and the subsequent analysis revealed that after 3 months the implant had degraded and been replaced by new bone tissue[128, 129]. At the end of this degradation process most of the alloying elements such as Al would have been released into the bodies of the rabbits. The long term health effects on the rabbits are unknown, but in the case of the human body, the release of Al into the body will create undesirable health problems[130]. In humans, Al is a neurotoxicant and its long term accumulation in brain t issues has been linked to neurological disorders such as Alzheimers disease, dementia and senile dementia[131]. In addition, the administration of RE elements such as cerium, praseodymium and yttrium has resulted in severe hepatotoxicity in rats[132]. Furthermore, using heavy metal elements as alloying components are also potentially toxic to the human body due to their ability to form stable complexes and disrupt the normal molecular functions of DNA, enzymes and proteins[133]. Therefore, there is a definite requirement to carefully select alloying elements that are non-toxic to the human body, see Table 4. Non-toxic alloying elements such as Ca[134] and Zr[135] have the potential to significantly improve the corrosion resistance of the Mg alloy and reduce the degradation rate to make the Mg metal alloy a viab le implant material[33].
Table 3. Chemical analysis of alloying elements for a selection of magnesium alloys
Alloy Nominal element component (wt. %) Maximum values of trace elements (wt. %)
Al Zn Mn Ca Li Nd Zr Y
AZ31 3.5 1.4 0.3 - - - - - Fe (max 0.003), Cu (0.008), Si (1.2), Ni (0.001) and Be (5 – 15 ppm)
AZ91 9.5 0.5 0.3 - - - - - Fe (max 0.004), Cu (0.025), Si (0.05), Ni (0.001) and Be (5 – 15 ppm)
AM60 6.0 0.2 0.2 - - - Fe (max 0.004), Cu (0.008), Si (0.05), Ni (0.001) and Be (5 – 15 ppm)
LAE442 4.0 4.0 2.0 Contains some heavy metal rare earth elements
WE43 - - - - - 3.2 0.5 4.0 Contains some heavy metal rare earth elements
Note: Table compiled from references[92, 93, 216, 217 and 218]
Gérrard Eddy Jai Poinern et al. : Biomedical Magnesium Alloys: A Review of Material Properties, 226 Surface Modifications and Potential as a Biodegradable Orthopaedic Implant
Table 4. Common alloying elements used in magnesium alloys
Alloying Element
Mechanical Properties Enhancement to Mg Matrix Pathophysiology Toxicology
Aluminium
Rapidly diffuses through Mg matrix, and acts as a
passivating element and improves corrosion
resistance. Improves die cast-ability
Blood serum level 2.1-4.8 µg/L
Tends to diffuse out of Mg matrix Neurotoxic (influences function of the
blood brain barrier) Linked to Alzheimer’s disease
Accumulates in amyloid fibres/brain plaques.
Accumulates in bone tissue/decreases osteclast viability
Calcium Adding to improve
corrosion resistance in Mg-Ca alloys.
Blood serum level 0.919-0.993 mg/L. - Levels controlled by Homeostatis of skeleton. Abundant mineral that is mainly stored in bones and teeth. Activator/stabilizer of enzymes.
Involved in blood clotting
Metabolic disorder of calcium levels results in the formation of excess calcium in the
kidneys (stones).
Copper
Can increase strength of Mg casts, however, it also
accelerates corrosion rate when exposed to a NaCl
medium.
Blood serum level 74-131 µmol/L Essential trace element
Excessive amounts of Cu have been linked to neuro-degenerative diseases.
Can produce cellular cytotoxicity.
Manganese
Adding to reduce the harmful effects of impurities
and improve corrosion resistance
Blood serum level <0.8 µg/L Essential trace element
Influences cellular functions/immune system/blood clotting/bone growth.
Influences metabolic cycle of lipids/amino acids and carbohydrates
Excessive amounts of Mn can produce neurological disorder. (manganism)
Lithium Improvement in corrosion resistance
Blood serum level 2-4ng/g Used in drugs to treat psychiatric
disorders
Overdose causes central nervous centre disorders, lung dysfunctions, impaired
kidney function.
Rare earth Elements
Improvement in corrosion resistance
Many rare earth elements have anticancerogenic properties and are used
in the treatment of cancer. Accumulate in the liver and bone
Zinc
Improves yield stress, Mg alloys containing Zn have an Elastic Modulus similar to
bone. The presence of Zn can
reduce hydrogen gas evolution during
bio-corrosion.
Blood serum level 12.4-17.4µmol/L Essential trace element
Essential to enzymes and immune system
In high concentrations is neurotoxic and can hinder bone development.
Note: Table compiled from references[122, 134, 213, 222, 223, 224, 225, 226 227 and 228]
4. Surface Modifications and Treatment Processes for Biomedical Mg Alloys
The high degradation rate of Mg and Mg alloy implants in the human physiological environment would result in the reduction of mechanical integrity of the implant before the bone tissues had sufficient time to heal[26]. There are two methods of reducing the degradation rate; the first, which was discussed in Section 3, involved alloy ing Mg with biocompatible elements that can resist the corrosion process. The second method is d iscussed in this section and involves the surface modification of the implant, through a treatment process that provides a resistive barrier against the body environment. An important factor that needs to be taken into account before any surface treatment is investigated is the healing or regenerative processes of bone and other associated body tissues. The healing process consists of three phases; inflammatory, reparat ive and remodelling.
The in itial inflammatory phase usually lasts between 3 to
7 days and this is the natural response of the body’s immune system to the presence of the biomedical device or implant. The reparative phase usually takes 3 to 4 months, during which t ime integration of the implant with the new and regenerated tissues takes place. The final remodelling phase, which is the longest phase, can take from several months to years to complete[136]. For Mg to be an effective bio-absorbable implant the degradation rate must be slow enough for the healing process to take place and the new tissues have sufficient time to provide their own structural support before the structural integrity of the implant is compromised. The minimum period for this to take place is at least 12 weeks[26]. Unfortunately, Mg alloys can completely degrade before the end of this timeframe and as a result there is a need to reduce the biodegradation rate. The bulk propert ies of Mg based alloys dictate its mechanical properties, but it is the surface properties that influence the interaction between the metal and the surrounding tissue environment of the body. As a consequence, surface
227 American Journal of Biomedical Engineering 2012, 2(6): 218-240
modifications and treatments can have a significant ro le to play in governing the degradation rate of the implant. To date, numerous surface modification techniques have been developed to change the surface characteristics of biomaterials. Many of these methods have been applied to modifying the surface p roperties of Mg bio-alloys. A brief overview of some of these surface modificat ion processes are presented in the following the four sections.
4.1. Mechanical Modifications to Induce Surface and Subsurface Properties
The surface structure of an implant is very important, since it is the init ial response of the surrounding tissues to the surface of the implant material that determines whether or not there is effective t issue-biomaterial integration. Studies of conventional types of permanent implant materials have shown that surface roughness can influence both cell morphology, cell growth and implant integration. In addition, modification of the surface topography by the physical placement of grooves, columns, pits and other depressions can influence cell orientation and attachment[137-139]. In the case of Ti alloys, surface modifications such as grooves, surface sand blasting and acid etching has revealed that grooved surface features provide superior cell attachment and promote greater cell pro liferation than roughen surfaces[140]. For Mg alloys, the influence of different mechanical p rocessing operations during fabricat ion has the potential to great ly influence surface and subsurface properties[141, 142].
Mechanical processing techniques involve operations such as rolling, shot peening, and milling. In the case of milling, at low cutting speeds, the surface formed by honed cutting tools tends to produce a rougher surface than those of sharp cutting tools. Also, during milling and similar metal chip removing processes, the exact effect on the underlining sub-surface is not fully understood[143], while chip removal from the surface during machining can direct ly influence the surface topography[144]. Besides machining techniques for chip removal, the use of rolling operations can also generate high passive forces acting normal to the surface, which can induce work hardening of the sub-surface. During the rolling operation the sub-surface grain structure is changed by the compressive stresses induced and the resultant micro -topography of the surface is significantly changed[145]. A recent study by Denka et al. has revealed a significant reduction in the corrosion rate (a factor of 100 was achieved in corrosion studies) of an Mg-Ca alloy that was deep-rolled, compared to the same alloy that was mach ined[146]. The presence of residual compressive stresses after ro lling also has the advantage of reducing micro-crack formation from pre-existing crack nucleation points within the substrate. The suppression of crack formation is also an important factor in improving the fatigue life cycle of a material being considered fo r b iomedical applications[146, 147].
The importance of surface and sub-surface treatments on
Mg alloy implants was recently investigated by Von Der Hoh et al.[148]. In their study three surface machining treatments were applied to an Mg-Ca (0.8 % wt calcium) alloy. The alloy was used to make three different geometric sample types. The first test sample was a machined 3 mm diameter s mooth cylinder, the second was like the first, except that it was sand blasted for 30 s using particles ranging in size from 300 to 400 µm and the final surface topography was a threaded cylinder. The smooth cylinders were machined with no fu rther surface treatment, so they retained the micro-surface topography produced by the cutting tool. After 6 months of in vivo implantation in adult New Zealand white rabbits, the smooth cylinders revealed good integration with the surrounding tissues and also had the least structural loss. The sand blasted cylinders had the greatest material loss with the init ial cylindrical shape completely consumed, while the threaded cylinders ranged between these two extremes. The results indicated that the smoother micro-topographic surface features of the cylinders were suitable for resorbable Mg alloys, while the test samples with the rougher surfaces promoted higher degradation rates. The results of this study clearly indicated that differences in surface roughness of the test samples could significantly influence the in vivo degradation rates. The study also highlighted the need for further investigation into the effects of different surface modifications on other biocompatible Mg alloys.
4.2. Physical and Chemical Modifications
From an engineering point of view, the most effective way to prevent corrosion is to coat the metal component with a protective barrier that effect ively isolates the metal from the surrounding environment. To be effective against corrosion, the protective coating must be uniform, well adhered and free from any imperfections such as pits, scratches and cracks. The major problem with Mg, as mentioned earlier, is its chemical reactivity when exposed to air or an aqueous environmental which results in the formation of an oxide/hydroxide layer over the metal surface. The presence of the oxide/hydroxide layer will have a detrimental effect on the ability of the coating to adhere to the metal surface and form a unifo rm protective layer. Therefore, surface cleaning and a suitable pre-treatment o f the metal surface is a crucial factor in achiev ing an effective surface coating.
4.2.1. Physical Vapour Deposition (PVD) & Chemical Vapour Deposition (CVD)
The PVD process involves the deposition of thin layers of metal and metal alloys from atoms or molecules from the vapour phase onto a substrate surface. During the process a metal or metal alloy is heated in vacuum chamber until it evaporates and then the subsequent vapour condenses onto the cooler substrate. This process has been successfully used on a variety of metals, but in the case of Mg there are a number o f problems to overcome. For example, in most PVD processes the substrate temperature range is usually between
Gérrard Eddy Jai Poinern et al. : Biomedical Magnesium Alloys: A Review of Material Properties, 228 Surface Modifications and Potential as a Biodegradable Orthopaedic Implant
400 and 550 oC, but in the case of Mg the substrate temperature must be kept below 180 o C for material stability reasons. The lower substrate temperature of Mg also influences the adhesive and corrosion resistance properties of the coating[149, 150]. The PVD process has successfully deposited binary alloys such as Mg-Ti, Mg-Zr and Mg-Mn along with other less biocompatible and toxic alloys. The subsequent corrosion studies have revealed that the binary-alloyed surface coating were capable of increasing the corrosion resistance of the various Mg alloys[151-153]. The chemical vapour deposition process has also been used to produce a variety of coating processes that can create a protective coating or modify the existing Mg alloy surface. During the deposition of a solid material from the vapour phase onto a (usually) heated substrate a chemical reaction over the surface takes place. This results in changes to the sub-surface of the substrate, which chemically modifies the surface properties. For example, the deposition of diamond like carbon (DLC) films on metallic implants can improve the surface properties of the implant, thus making it biocompatible with the surrounding body tissues[154].
4.2.2. Ion Implantation and Plating
Ion implantation consists of bombarding the surface of a substrate with ionized particles. The ionized part icles penetrate the surface and become embedded in the sub-surface of the substrate. The ionized particles soon neutralize in the interstitial positions within the grain structure forming a solid solution. During this process physiochemical changes take place in the sub-surface of the substrate, while the bulk properties of the substrate remain unchanged. To date there have been relatively few studies carried out that have used ion implantation to enhance the surface properties of Mg alloys. A recent study by Liu et al. examined the corrosion behaviour of surgical AZ91 after it was subjected to Ti ion implantation[155]. The study revealed that a compact surface oxide layer was formed, which was predominantly composed of TiO2 with a smaller amount of MgO. Subsequent testing in simulated body flu id at 37 ± 1 revealed that the corrosion resistance of ion treated AZ91 alloy was improved significantly. In a similar study by Fang et al. the corrosion behaviour of a new medical grade Mg-Ca alloy was examined before and after Zn ion implantation[156]. The results revealed that after ion implantation, the Zn had improved the surface hardness and elastic modulus of the alloy. The surface oxide layer fo rmed during corrosion testing in simulated body fluid enhanced the Mg-Ca alloys corrosion resistance. However, Wan et al. examined a Mg-Ca alloy before and after Zn ion implantation, the results found that the ion-implanted substrates had a lower corrosion resistance than the untreated substrates[157]. Subsequent analysis of the results suggests that Zn was an unsuitable metal for ion implantation with Mg-Ca alloys for b iomedical applicat ions.
Ion plating is a technique that deposits noble metal ions onto a less noble metal substrate to form a dense and
well-adhered layer. The plating layers improve surface properties such as topography, roughness, surface chemistry and wear resistance. Zhang et al. used this technique to plate a pure Mg substrate with Ti ions and then subsequently studied its corrosion behaviour in a 0.9 wt % NaCl solution[158]. The results not only revealed a substantial improvement in corrosion resistance, but also found that an interfusion layer had formed between the Ti coating and the Mg substrate.
4.2.3. Thermal Spray Coatings
During th is coating process, materials such as metals, metal alloys, ceramics, polymers and composites are feed (powder or wire form) into a gun. The material is then heated to a molten or semi molten state within a gas stream. The resulting micrometer size droplets are accelerated in the gas stream, which is directed towards the surface of the substrate[159]. This technique was successfully used by Zhang et al. to deposit an Al layer on an AZ91D substrate[160]. To ensure adhesion of the coating to the substrate, a post heat treatment process was carried out 450 . There was significant diffusion of Al and Mg around the interface of the coating, which enhanced both the corrosion resistance and anti-wear properties of the coating. The disadvantage of using an Al coating on the AZ91D substrate for a possible biomedical implant applicat ion is the negative effect of Al3+ ions being released into the surrounding tissues during subsequent corrosion. Ceramic coatings such hydroxyapatite (HAP), TiO2, Al2O3, and ZrO2 have also been successfully applied to Ti alloys to improve their corrosion resistance, wear resistance and biocompatibility[161]. However, in a study by Zeng et al., thermally sprayed TiO2 onto an Mg alloy (AM60) revealed that the subsequent coating showed no improve in its corrosion resistance compared to the untreated Mg alloy when they were both immersed in Hanks’ solution[162]. The study also revealed that galvanic corrosion occurred between the surface of the Mg alloy and the coating layer, which effectively reduced any protective properties offered by the coating. This highlights the weakness of thermally sprayed ceramic coatings which have rough surfaces, high porosity and poor adhesion properties[163].
4.2.4. Laser Surface Melting, Alloying and Cladding
The high-density energy of a laser beam can be effectively used to modify the surface region of Mg alloys. The surface region of the alloy can be melted to create a meta-stable solid solution. This is then followed by rapidly cooling the substrate, which results in the refinement of the surface microstructure. This technique can also be used to improve the surface properties of an Mg alloy substrate by melting a metallic coating and the underlin ing sub-surface. During the rapid melting process both the coating and sub-surface mix before re-solidifying during subsequent cooling to form a new surface alloy which coats the bulk of the substrate. Furthermore, if the appropriate alloying metals are
229 American Journal of Biomedical Engineering 2012, 2(6): 218-240
incorporated into this surface modificat ion technique it is possible to significantly improve surface properties such as corrosion resistance[164, 165]. For example, improved surface properties of a Mg alloy (AZ91) have been achieved with the dispersion of hard metallic part icles such as TiC and SiC in the molten pool generated by laser melting[166, 167]. Also, laser cladding of an Al-Si alloy onto a number of Mg alloys such as AS41, AZ91 and W E54 have also been attempted, but unfortunately, the surface properties were not significantly improved[168, 169].
4.3. Wet Chemical Processes
4.3.1. Electrochemical Deposition of Metallic Coatings
The corrosion resistance of Mg and its alloys can be increased by an electroplating technique. In this technique a metal salt is reduced in solution to its metallic form, the electrons for reduction are supplied from an external source and the resulting metallic ions are deposited on to the surface of the substrate. However, most metals are more electrochemically noble than Mg, which can cause serious problems if there are any imperfect ions in the deposited layer. Such imperfections will expose the underlin ing substrate and result in the format ion of s mall localized areas of corrosion. The corrosion sites form h ighly corrosive pits that tunnel down into the Mg substrate and seriously weaken the substrate[141]. From an industrial point of view, electroplating is a highly effect ive technique for coating Mg and its alloys with metallic coatings such as nickel, ch rome and alumin ium coatings[142]. These coatings have good mechanical properties and provide effective corrosion protection. Unfortunately, these metals are also harmful to human t issues, which make them h ighly unsuitable for biomedical applicat ions.
4.3.2. Chemical Conversion Coatings
Chemical conversion coatings are formed by chemically treating the surface of Mg and Mg alloys to produce a thin outer coating of metal oxides, phosphates or other compounds that are chemically bonded to the surface[170]. The conversion coating acts as protective barrier that isolates the substrate from the surrounding environment and prevents the corrosion.
Industrially, there are several different types of conversion coatings such as chromate, phosphate/permanganate, rare earth, stannate and hydrides. Many of the processes used to produce conversion coatings involve the use of toxic materials that are detrimental to human health. For example, the presence of hexavalent chromium (Cr6+), that is used in chromate coatings.
An alternative treatment to chromate conversion coatings are: 1) phosphate; 2) phosphate-permanganate; and 3) stannate coatings. All three of these conversion treatments have comparable corrosion resistant properties to those of chromate t reatments. Xu et al. have investigated the behaviour of a phosphate treatment on an Mg alloy
(Mg-Mn-Zn) that was subsequently immerged in a simulated body solution (SBF). During the treatment process a biocompatible brushite layer[CaHPO4
.2H2O] was formed on the surface of the substrate. Subsequent immersion in SBF revealed that the brushite layer transformed into a coating of HAP, with the excess phosphate ions being released into the surrounding environment. The released phosphate ions were also found to neutralize the alkalizat ion effect produced by the corrosion process. The treatment process did not prevent corrosion, but it did significantly slow down the degradation rate[171]. And a phosphate-permanganate process developed by Han et al. using Mn3(PO4)2 was ab le to produce a resilient surface coating on a Mg alloy (AZ31D) that was self healing in saline solutions[172]. While a stannate treatment developed by Gonzalez-Nunez et al. was able to deposit a 2 to 3 µm thick layer of MgSnO3 on an Mg alloy (ZC71). The coating was adherent, continuous, and crystalline which produced a passivating effect on the substrate surface[173]. Unfortunately, no degradation rate data was reported, indicating that more studies are needed to indeed gauge the effectiveness of this process for in vivo applications.
Magnesium fluoride (MgF2) conversion coatings on Mg alloys have produced mixed results in providing corrosion protection. Degradation studies carried out by Zeng et al. reported that an MgF2 conversion layer formed on an Mg alloy (AZ31) provided marg inal corrosion resistance in a 0.9 wt % NaCl solution[174]. While in a similar study Hassel et al. found that a conversion layer of MgF2 formed on an Mg alloy (ZM21) could provide reasonable corrosion resistance[175]. And an in vivo study carried out by Witte et al. revealed that a conversion layer thickness between 150 and 200 µm formed on the surface of a RE based Mg alloy (LAE442) was able to reduce the degradation rate and reduce the release of alloying elements[176]. In a similar study by Gao et al., the feasibility of forming RE conversion layers on pure Mg to improve corrosion resistance was examined[177]. The two RE elements under investigation were Ce and Y, with each element being used individually to form a surface treatment solution. The first solution contained CeCl3 which formed a conversion layer consisting of Mg (OH)2, CeO3 and MgO, while the second solution contained Y(NO3)2 which formed a conversion layer consisting of Mg (OH)2, Y2O3 and MgO. The study revealed that the second conversion layer had improved corrosion resistance compared to the first. However, both coating provided limited corrosion resistance due to their thin thickness and soft structure, which was incapable of withstanding minor mechanical damage. Furthermore, both the toxico logy and metabolic pathways within the human body of RE elements such as Ce and Y are still unclear and need to be fully investigated before they can be used in biomedical applications.
4.3.3. Calcium Phosphate Surface Coatings
A more biocompatible form of conversion coating can be derived from a variety of calcium phosphate compounds. In
Gérrard Eddy Jai Poinern et al. : Biomedical Magnesium Alloys: A Review of Material Properties, 230 Surface Modifications and Potential as a Biodegradable Orthopaedic Implant
particular, HAP[Ca (PO4)6(OH)2], which has been widely used as a bone substitute and replacement in several biomedical applicat ions[178-180]. There are three major advantages in using HAP in hard tissue engineering applications: 1) it has good biocompatibility and bioactivity properties with respect to bone cells and other body tissues; 2) it has a slow biodegradability in situ; and 3) it offers good osteoconductivity and osteoinductivity capabilit ies[181, 182]. These properties are very important because bone tissue constantly undergoes remodelling, a process in which bone tissue is simultaneously replaced and removed by the bone cells, (osteoblasts and osteoclasts respectively). It is these advantages that make HAP and TCP (tri-calcium phosphate) compounds attractive for coating metallic orthopaedic implants. In this application, both HAP and TCP coatings promote bone formation which enhances bonding between the implant and the surrounding tissues.
It is also due to these positive biolog ical responses within the human body that has made calcium phosphate coating an attractive option for potentially reducing the biodegradation rate of Mg orthopaedic implants. Several techniques have been used to deposit calcium phosphate coatings onto Mg substrates, these range from anodization[183], b io-mimet ic co at ings[1 84- 18 6], ele ct ro- de posi t ion[1 87, 18 8], hydrothermal[189] and wet chemical methods[190, 191].
Xu et al. has investigated using an immersion technique that involves soaking an Mg-Mn-Zn alloy in an alkaline solution to form a b rushite (CaHPO4.2H2O) surface coating. Unfortunately, the layer formed was porous and did not prevent corrosion in a simulated body fluid. However, the degradation rate was significantly reduced and provided the Mg alloy substrate with reasonable protection against the corrosive effects of the simulated body flu id. The study also found that the brushite coating was able to improve the surface biocompatibility of Mg alloy substrate, since the brushite coating transformed into a HAP phase with time. Also during this transformation, acid ic phosphate ions were released into solution, which tended to have a neutralizing effect on the alkalizat ion process[192]. Furthermore, the surface treatment enhanced the bioactivity of the Mg-Mn-Zn alloy and promoted bone format ion[193].
A similar calcium-phosphate coating was produced by Wang et al., which involved immersing a Mg substrate into a solution containing Ca and P[Ca (NO3)2 and Na2HPO4] to create a di-calcium phosphate di-hydrate (DCPD) surface layer[194]. The DCPD layer was effective in providing protection for the Mg substrate during the first 21 days of immersion in a simulated body fluid.
Recently, Yanovska et al. investigated the influence of low magnetic fields during a one-step dipping technique [195]. The involved dipping a Mg substrate into an aqueous solution containing Ca (NO3)2.4H2O and Na2HPO4.12H2O. Deposition of both DCPD and HAP phases under the influence of magnetic fields lead to crystal orientation during the formation of the phases. The technique also produced coatings with enhanced corrosion resistance, which in turn reduced the degradation rate[195]. In an alternative method,
Song et al. used an electro-deposition technique to produce three types of coating namely; DCPD, HAP and fluorapatite (FHA). The study found the FHA coating had long-term stability and remained intact even after 1 month of immersion in a simulated body fluid and provided effective corrosion resistance to the Mg alloy[196].
4.3.4. Alkali Heat Treatments
Heat treatment can be a beneficial way of improving the microstructure and enhance the surface properties of Mg and Mg alloys. The corrosion behaviour and cytotoxicity of alkali heat-treated pure Mg samples immersed in simulated body fluid (SBF) were investigated by Li et al.[197]. The samples for treatment were p laced into a super saturated NaHCO3-MgCO3 solution and then heat-treated. SBF solutions with and without chloride ions were used to study the influence of chloride ions on the corrosion behaviour of treated Mg samples. A ll the treated samples showed a significant improvement in corrosion resistance in both SBF (Cl-) and SBF solutions compared to the untreated Mg samples. In addition, after 14 days of immersion in the SBF flu ids, a calcium phosphate compound with a molar ratio of 1.858 was detected on the surface of the samples. While the subsequent cytotoxicity testing revealed no signs of morphological changes in the cells and no inhibitory effect of the surface treatments on cell growth could be detected.
In a similar study by Liu et al., the corrosion behaviour of a heat-treated Mg-Al alloy (AZ63) immersed in a SBF solution for 14 days was investigated[198]. During heat treatment (solution treatment at 413 o C for 24 h followed by aging at 216 for intervals of 1 h, 5.5h and 12 h), the microstructure of the Mg-Al alloy changes as the Al atoms in the alloy diffuse towards the grain boundaries and subsequently precipitate out of solution to fo rm the β phase. And as a consequence of the diffusion process (Aging), the concentration of Al atoms remain ing in the matrix (α phase) decreases and results in the matrix having a reduced corrosion rate. In addition, the study found that samples microstructure significantly influenced the overall corrosion morphology. For example, the surface of the untreated samples displayed deep and uniform corrosion, while the surface of the treated samples had only shallow pitting[198].
4.3.5. Anodization
The anodization of magnesium is an electro-chemical process that changes the surface chemistry of the metal, v ia oxidation, to produce a stable anodic oxide layer. The structure of this layer is characterized by a thin barrier layer at the metal-oxide interface, fo llowed by a less dense porous oxide layer. The porous layer can d isplay a variety of different structures and properties which are dependent on the composition, substrate micro-structure and processing parameters[142]. The processing parameters that influence oxide layer formation include: 1) the type, temperature and concentration of electrolyte; 2) current density; and 3) the applied anodization voltage. These parameters can also
231 American Journal of Biomedical Engineering 2012, 2(6): 218-240
significantly influence the resulting corrosion behaviour of the substrate[199]. The anodization process can also produce an oxide layer consisting of pores, whose size and density is dependent on the selection of the appropriate processing parameters. Industrially, porous oxide layers are usually coloured and then sealed or form part of a pre-treatment process prior to painting or coating. Many of the industrial coating and surface treatments used on anodized Mg alloy components to reduce corrosion are toxic to the human body. And as a result, research efforts have focused on searching for biocompatib le surface treatments and process that are non toxic. For example, Hiromoto et al. have studied the effects of controlled calcium phosphate (Ca-P) precipitation on pure Mg substrates by anodization and then thermally treating the substrates in an autoclave[200]. After thermal treatment, the subsequent immersion studies in Hanks’ solution revealed that the Ca-P coated substrates had very litt le corrosion[201]. The advantage of this technique comes from the Ca and P elements being deposited on the substrate, since both bioactive materials are known to induce osteoinduction and promote new bone tissue growth[201].
Micro-arc oxidation (MAO) is an electrochemical process which uses a high anodic voltage and high current density to create an intense micro-arc (plas ma) near the metal surface to induce oxidation. The oxide layer formed during this process is substantially thicker than conventional anodization, since the sub-surface of the metal substrate is also oxidized[202]. This technique can be used to deposit ceramic coatings on valve metals such as Al, Mg, Ta, Ti, W, Zn and Zr and their alloys. And by selecting the appropriate process parameters, the (MAO) technique can produce high quality coatings with superior adhesion, corrosion resistance, micro -hardness, wear resistance and strength. In a corrosion and wear study by Zhang et al., die cast Mg alloy (AZ91D) substrates were treated with a MAO coating. Then both treated and untreated substrates were immersed in the Hanks’ solution to determine the effectiveness of the MAO coating in reducing the corrosion rate[203]. Immersion testing revealed that the untreated substrates loss 15 times more mass due to corrosion than the MAO treated substrates. In addition, the mass loss from the untreated substrates during wear tests was 1.5 times greater than those of the MAO treated substrates. The tests clearly indicated that the MAO surface treatment was effect ive in improving both the corrosion and wear resistance of the Mg alloy. Unfortunately, there is no current data available describing the combined effects of corrosive and wear on MAO treated Mg alloys in Hanks solution.
In a recent study examin ing the combined effects of corrosion and wear on an Mg Alloy (AZ91), Chen et al., investigated untreated and MAO treated substrates immersed in different composition based solutions composed of NaCl and NaHCO3[204]. The study revealed that the MAO coating did improve the corrosion resistance of the Mg alloy. However, the wear resistance of the treated substrate tended to decrease with t ime. The decreasing wear resistance was
caused by wear debris, which consisted of abrasive particles produced by the breakdown of the MAO coating. Both anodization and MAO oxide coating can effect ively reduce the corrosion rate of Mg and Mg alloys by producing a strong resilient oxide layer that provides an effective protective barrier between the metal substrate and the fluid environment. For example, Song et al. has demonstrated that the oxide layer produced during the anodizat ion of a pure Mg substrate was capable of providing an effect ive barrier to corrosion for 1 month in Hanks’ solution. During this time no hydrogen evolution was detected, indicating that the degradation had been delayed[205, 206]. This is an important factor for a biodegradable implant, since controlling the dissolution rate allows matching of the degradation rate of the implant with the growth o f new replacement bone tissue.
4.4. Polymer Coatings
Many implantable biomedical devices and implants are coated with a thin adherent polymeric material that effectively isolates the device from the flu idic environment of the body. Polymer coatings are frequently used to modify the surface properties of biomedical implants to improve their biocompatibility, performance and therapeutic effectiveness. The interface between implant surface and body environment is critical in soliciting the appropriate immunological response. Therefore, selecting the correct polymer coating is crucial in determin ing the biocompatibility of the implant and also provides a wider range of design options that can be used to improve the surface properties of the original implant surface. For example, selecting a polymer coating which slowly biodegrades can potentially delay the corrosion of an Mg implant and maintain its mechanical integrity over a longer timeframe.
Biomedical coatings can be divided into two primary categories: 1) short term, which include d isposable or single patient use; and 2) long term applicat ion of prosthetic implants and reusable laboratory equipment[207]. To achieve the bio-functional requirements and protection, a successful polymer coating must adhere to the Mg implant and be strong and flexib le enough to withstand the normal movement of the implant. The coating should also be capable of being sterilised and be sufficiently durable to perform its protective function under the expected conditions of the particular application[208]. The advantage of polymer coatings is that they can be chemically, physically and mechanical customized to suit a specific application. Thus, being able to select and fine-tune the various properties has enabled polymeric materials to be used in a wide range of coating applications such as protection, improved lubricity, antimicrobial, adhesion resistance, ultrasonic imaging and blood compatible coatings for drug delivery[209].
Before coating, an appropriate pre-treatment process is required to effectively clean the surface of the Mg implant. This process should produce a clean, dry and contaminant free surface capable o f provid ing the maximum possible
Gérrard Eddy Jai Poinern et al. : Biomedical Magnesium Alloys: A Review of Material Properties, 232 Surface Modifications and Potential as a Biodegradable Orthopaedic Implant
adhesive strength between the polymer and implant surface. The presence of entrapped air and moisture on the surface could lead to degassing and the format ion of holes in the coating during the curing process. The coating process and associated coating parameters can influence the resulting microstructure and morphology of the polymer coating. This explains why coatings with similar compositions can have different surface properties[210]. A polymer coating can be applied to an Mg implant using a solvent, an aqueous solution or from the vapour phase. Water based coating techniques have the advantages of eliminating or reducing solvent effects on the substrate and are more environmental friendly, since there are no toxic solvent wastes produced. Polymer coating produced by vapour-deposition can produce both uniform and defect-free coatings. The quality and chemical structure of these coatings combine to provide an effective barrier to the body environment and also enhance the protective properties of the coating. For example, Parylene (poly (p-xy lylene), is a polymer coating that has been used to protect implanted sensors and other biomedical devices[210].
In addition, Mg needs a suitable surface primer or b inding agent to improve the adhesion between the polymer coating and the implant surface. Unfortunately, primers composed of metallic ions, components of binding agents and almost all organic solvents found in paints and similar surface preparations are toxic or detrimental to the human body and therefore cannot be used in biomedical applicat ions. Recently, Huang et al. primed a pure Mg substrate using a silane coupling agent before using dip coating technology to coat the substrate with degradable poly (lactic acid)[211]. The silane coupling agent was found to improve the adhesion between the poly (lactic acid) and the substrate. Similarly, PLGA, a polymer with good blood compatibility was also coated onto a pure Mg substrate using the same process. Subsequent corrosion testing in Hanks solution revealed that the PLGA could provide an effective coating and protect the underlining Mg substrate. However, corrosion protection of the Mg substrate was dependent the degradation of the PLGA. And in a bone replacement application, the coating was unable to provide an adequate protective barrier[211]. In a similar study, Xu et al. was able to show that producing different surface modifications on chitosan coatings applied to Mg alloy substrates were able to influence the biodegradability of the coating[212].
Polymeric coatings have the potential to modify the surface properties of an Mg based implant and significantly improve the implants usability, durability and performance. Polymeric materials used in coating an implant have the potential to be specifically designed to provide physical, chemical and mechanical responses that can be fine-tuned to solicit and enhance specific biochemical responses within the body environment. However, a significant research effort is still needed to find polymeric materials that can produce thin, dependable, mult i-functional coatings with controllable degradation rates that can be used to prolong the mechanical effectiveness of Mg based implants.
5. Conclusions
Mg and Mg based alloys are ext remely biocompatib le and have similar mechanical properties to natural bone. This makes them an attractive material for the manufacture of biodegradable, with the capability to replace many currently used orthopaedic materials such as biodegradable biopolymers. And despite having the potential to function as an osteoconductive and biodegradable substitute in load bearing applicat ions, the practical application of Mg based alloys faces the serious challenge of overcoming the rapid corrosion rates that occur within the physiological environment of the body. The types of biological corrosion occurring within the body environment and the influence of body fluid pH, concentration of ions, protein adsorption on the implant surface and the influence of the surrounding tissues was discussed. To overcome the effects of b iological corrosion, a number of treatment methods designed to reduce the corrosion rate, such as the addition of alloying elements and surface modification techniques were discussed. The biological consequences of adding alloying elements to reduce the corrosion rate was explored and the need for careful selection was d iscussed. For example, alloying elements such as Al and Li can be used to improve the corrosion resistance of an alloy, but the release of their ions in the body can create undesirable health problems. Therefore, careful selection of the appropriate alloying elements and the resulting corrosion by-products were discussed in this article. Both these issues are critical for any material being considered for a b iomedical application within the human body.
Furthermore, the development of new b ioactive surface modifications and coatings, with superior physiochemical and mechanical properties has the potential to enhance the performance of Mg alloy implants by improving both their corrosion and wear resistance. The first type of surface modification discussed used conventional mechanical processes such as machining and rolling to enhance both surface and sub-surface properties. The second type of surface modification discussed examined the types of physical and chemical surface treatments that could deposit a metallic or ceramic coating or produce a surface conversion coating. The corrosion resistance and biocompatibility of the various surface modifications were discussed. In the case of ion implantation, the technique gave mixed corrosion resistance values when Zn was used to treat an Mg-Ca alloy. While a conversion coating technique using a calcium phosphate compound produced a coating that not only reduced the corrosion rate, but also improved the biocompatibility and promoted bone format ion at the surface of the Mg alloy. Whilst the use of polymeric coatings to protect implanted sensors and biomedical devices is well established, further studies are needed to examine the viability of using a polymeric material with suitable biodegradable properties to extend the operational life of a Mg alloy implant.
The combination of new Mg alloys and evolving surface
233 American Journal of Biomedical Engineering 2012, 2(6): 218-240
modification processes, has presented the opportunity to design and develop a biocompatib le material that has the potential to be used in an orthopaedic implant. The ability to select alloying elements and surface modifications provides the opportunity to design a specific Mg alloy implant with mechanical properties and biodegradation profile that can be tailored to the specific orthopaedic application. However, more biomedical studies are needed to investigate the interaction between the material surface and the surrounding tissue environment. Also, more in situ experimental studies are needed to examine the long-term effects of alloying elements released during the bio logical corrosion of Mg based alloys. The biomedical, materials science and engineering research presented in this review article has clearly demonstrated the potential o f using Mg based alloys to manufacture orthopaedic implants, but there are still challenges to be overcome. The first challenge is to improve the corrosion resistance of Mg base alloys by using only biocompatible alloying elements and an appropriate biocompatible surface treatment that permits the controlled degradation of the implant. While the biomedical challenge consists of more clinical trials to establish the long-term biocompatibility of Mg based alloys and their corrosion products within the body environment.
ACKNOWLEDGEMENTS Dr Derek Fawcett would like to thank the Bill & Melinda
Gates Foundation for their research fellowship.
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International Journal of Biomedical Materials Research 2014; 2(2): 7-14
Published online September 30, 2014 (http://www.sciencepublishinggroup.com/j/ijbmr)
doi: 10.11648/j.ijbmr.20140202.11
ISSN: 2330-7560 (Print); ISSN: 2330-7579 (Online)
Chemical immersion coatings to improve biological degradability of magnesium substrates for potential orthopaedic applications
Sridevi Brundavanam, Gérrard Eddy Jai Poinern, Derek Fawcett
Murdoch Applied Nanotechnology Research Group, Department of Physics, Energy Studies and Nanotechnology, School of Engineering and
Energy, Murdoch University, Murdoch, Western Australia 6150, Australia
Email address: [email protected] (G. E. J. Poinern)
To cite this article: Sridevi Brundavanam, Gérrard Eddy Jai Poinern, Derek Fawcett. Chemical Immersion Coatings to Improve Biological Degradability of
Magnesium Substrates for Potential Orthopaedic Applications. International Journal of Biomedical Materials Research.
Vol. 2, No. 2, 2014, pp. 7-14. doi: 10.11648/j.ijbmr.20140202.11
Abstract: Historically, cobalt-chromium, stainless steel and titanium alloys have been the main principal materials used in a
variety of medical procedures for load-bearing implants in the body. Magnesium and magnesium-based alloys have the potential
to be used as short-term structural support during the healing process of damaged hard tissues and diseased bone. Unlike
traditional biologically compatible metals, which are not biologically degradable, magnesium based alloys offer both biological
degradability and biological absorbability. Despite the many advantages offered by magnesium, its rapid degradation rate in the
highly aggressive and corrosive body fluid environment has severely limited its present day medical application. This article
reviews the chemical immersion technique for producing calcium phosphate coatings on magnesium substrates for slowing down
the degradation rate while maintaining the biological compatibility and absorbability.
Keywords: Magnesium, Biodegradability, Chemical Immersion, Bone Tissue Engineering
1. Introduction
Biologically compatible materials have been engineered
into a variety of medical devices and implants to assist in
healing, replace diseased tissues and in some drastic cases
completely replace damaged tissues. The repair and
replacement of diseased or damaged hard tissues presents a
major challenge to patient health, wellbeing and quality of life.
A major consequence of tissue damage or loss resulting from
accidents or diseases is the psychological impact on patient
wellbeing. For example, even minor injuries to fingers or toes
that interfere with function and usually heal without much
trouble can have a significant impact. While severe bone
diseases and injuries that result in the loss of a limb creates
functional problems for patients for many years. In many
cases the device or implant is temporary and only needs to
remain in the body during the healing process. Once the
healing process is complete, a second surgical procedure is
required to remove the device or implant which significantly
increases patient site morbidity and associated health costs
[1-3]. Alternatively, when an implant needs to remain in the
body permanently, as in the case of a total joint replacement,
long-term biocompatibility, mechanical strength and
structural stability become important factors that must be
addressed.
Biologically compatible polymers have been extensively
investigated since the 1950’s for a variety of potential tissue
engineering applications. Natural polymers such as collagen
[4, 5], chitosan [6, 7], hyaluronic based derivatives [8, 9],
polysaccharides [10, 11], and a variety of protein based
materials such as fibrin gel[12, 13] have all been extensively
studied and found to be suitable for a wide range of tissue
engineering applications. Synthetically manufactured
biologically degradable and biologically absorbable polymers
such as Poly (lactic acid), PLA [14, 15], Poly (L-lactic acid),
PLLA [16, 17], Poly (lactic-co-glycolic acid), PLGA [18, 19],
Poly-caprolactone PCL [20, 21] and Poly (glycolic acid) PGA
[22, 23] have all being investigated and used in a variety of
biomedical applications. For example, biodegradable sutures
currently in clinical use are made from PLA and PGA. These
have also been extensively investigated for the controlled
delivery of drugs to specific organs within the body [24-26].
Advantages of using biological compatible polymers arise
from their low toxicity within the body and the ability to
8 Sridevi Brundavanam et al.: Chemical Immersion Coatings to Improve Biological Degradability of Magnesium
Substrates for Potential Orthopaedic Applications
control their degradation rate. Furthermore, the by-products of
degradation can be easily handled by the body’s natural
processes and excreted in the urine [27]. Polymers can also be
produced in a variety of shapes and structures such as disks,
films, fibres and pellets to meet the specific requirements for a
particular application. In addition, polymers can be produced
with micrometre and nanometre scale typographical surface
features to enhance cell-substrate interactions with the surface
of the implant [28].
Unfortunately, polymers with all their many advantages are
limited by their low mechanical strength, which severely
restricts their use in load bearing and hard tissue supporting
applications. Metals have more desirable mechanical
properties due to their relatively high strength, elastic modulus,
fracture toughness and resilience and as a result several
metallic biomaterials such as cobalt-chromium-based steel
alloys, titanium-based alloys, nickel-based alloys and stainless
steels have been widely used as implant materials [29-31].
However, studies have shown that conventional surgical metal
alloys are not biologically absorbable and because of
corrosion and wear, there is a release toxic metallic ions into
surrounding cells and tissues [32-34]. These detrimental
metallic ions induce an unfavourable inflammatory response
from the body’s immune system and the surrounding tissues,
which significantly reduces the biocompatibility of the
implant [33]. Furthermore, the significant difference in
mechanical properties between metal implants and
surrounding bone tissue results in a clinical phenomenon
known as stress shielding. For example, the elastic modulus of
both cobalt-chrome alloys and stainless steel is approximately
ten times larger than that of bone, while a titanium alloy such
as Ti-6Al-4V is around five times greater [35]. Normally, bone
tissues are constantly undergoing remodelling and
modification in response to the stresses produced during
everyday activities. However, the presence of a metal implant
creates stress-shielding, which results in a major portion of the
load been carried by the implant. Thus, with most of the load
been carried by the implant, the surrounding bone tissues
experience significantly less load related stress and as a result
leads to bone resorption, mechanical instability and the
ultimate failure of the implant [36]. In addition, metallic
implants used as temporary structural supports, such as pins,
screws, and plates often need to be removed by a second
surgical procedure once the healing process has taken place.
The increased health costs and morbidity associated with the
second surgical procedure highlights the need for new
biologically compatible materials that can provide short-term
structural support during the healing process. Then after
healing has taken place to an acceptable level, the material
would then biologically degrade and safely be reabsorbed and
metabolized by the body.
One interesting alternative to conventional metals used as
current bio-implants is magnesium. Magnesium (Mg) is a
lightweight, silvery-white metal that has been extensively
used in alloy form in a wide range of engineering applications
such as aerospace and automotive [37]. The density of Mg and
its alloys are around 1.74 g/cm3 at 20
oC, which is 1.6 and 4.5
times less dense than aluminium and steel, respectively.
Interestingly, the density of pure Mg is 1.74 g/cm3, while
natural bone ranges from around 1.8 to 2.1 g/cm3 and the
elastic modulus of Mg and human bone are 45 GPa and 40 to
57 GPa respectively [38, 39]. It is because of the close
similarity in the respective densities and elastic moduli that
have made Mg a promising candidate for hard tissue
engineering applications. The mechanical properties of Mg
being similar to natural bone means that it has the potential to
significantly reduce the possibility of stress shielding and
prevent the associated bone resorption problems. Mg is also
biologically degradable and biologically absorbable, with both
Mg and its corrosions products considered physiologically
beneficial, with as much as 30 g stored in the bone tissues and
muscles of an adult body [40]. The body uses Mg, a bivalent
ion, in a number of metabolic processes and to form apatite in
the bone matrix [41]. And recent studies by Robinson et al.
have shown that Mg has novel antibacterial properties against
pathogens such as Escherichia coli, Pseudomonas aeruginosa
and Staphylococcus aureus [42]. Because of these
advantageous properties Mg has gained significant interest as
a potential biologically degradable material that removes the
need for additional surgeries to reclaim pins, screws, and
plates used in the short term while the healing process takes
place. Figure 1 presents an ideal life span of an Mg implant
that slowly degrades and allows regenerating bone tissues to
progressively carry the load. The use of Mg as an implant
material also has the potential to avoid the long-term
complications associated with conventional metal implants in
the body.
Figure 1. Graphical representation of an ideal load carrying transition
between a slowly degrading Mg implant and progressively regenerating bone
tissue
2. Biological Degradation of Magnesium
The main limitation that prevents Mg being used in
orthopaedic applications is its low corrosion resistance in
body fluids, which are composed of water, dissolved oxygen,
proteins and electrolytic ions such as chloride and hydroxide.
In this highly corrosive aqueous environment results in the
rapid release of ions from the metal surface which combine
International Journal of Biomedical Materials Research 2014; 2(2): 7-14 9
with ions in the fluid to form chemical species, such as metal
oxides, hydroxides, chlorides and other compounds [43].
Generally metals have a tendency to corrode in electrolytic
environments with some metals having a greater propensity to
degrade more rapidly than others as seen in Table 1.
However, the interfacial region has also a significant bearing
on the overall performance of metallic bio-implants. This can
be shown in table The formation of metal oxides results in the
creation of an oxide layer composed of Mg(OH)2 that adheres
to the metal surface. The oxide layer is slightly soluble and
reacts with chorine ions to form highly soluble magnesium
chloride and the rapid production of hydrogen gas bubbles [44,
45]. The localised formation of gas bubbles generally begins
just after surgery and continues for periods as long as three
weeks. During this post surgery period, the pH around the
implant increases and results in alkalization of the surrounding
tissue environment. The presence of hydrogen bubbles and
local alkalization can severely affect pH dependent
physiological processes in the vicinity of the implant and
delay tissue healing [46]. However, the use of a subcutaneous
needle can be used to prevent significant build up of gas
around the implant. Furthermore, in his study Song has
suggested that small hydrogen evolution rates of around 0.01
ml/cm2/day can be easily handled by the human body and does
not constitute a serious threat [47]. The high initial corrosion
rate produces a thick oxide layer, which fully covers and seals
the metal surface to form a passive layer, which physically
stops or severely limits the migration of ionic species and
hydrogen gas across the metal oxide solution interface.
Unfortunately, the high corrosion rate during the first two to
three weeks can cause a significant reduction in the
mechanical and structural integrity of the implant before the
bone tissues have had sufficient time to fully heal.
Table 1. Electrochemical series of selected metallic ions and their voltage
potential [78]
Metal ions Potential (Volts)
Au3+ +1.420
Pt2+ +1.200
Cu2+ 0.345
Cd2+ -0.402
Fe2+ -0.440
Cr2+ -0.710
Zn2+ -0.762
Al3+ -1.670
Mg2+ -2.340
In atmospheric air at room temperature Mg corrodes to
form a thin grey oxide layer over its surface. The oxide layer
then reacts with atmospheric moisture to form a more stable
magnesium hydroxide [Mg(OH2)] and hydrogen gas
[43].Under standard environmental conditions, the Mg(OH2)
layer is able to provide some degree of protection and is also
capable of slowing down the corrosion rate even under
aqueous conditions [48]. When Mg is exposed to an aqueous
environment the corrosion process can be expressed by the
following equations. The primary anodic reaction involves
metallic Mg being converted to Mg2+
ions as seen in equation
(1), meanwhile the reaction occurring at the cathode,
presented in equation 2, involves the reduction of protons.
Anodic reaction:
Mg → Mg2- + 2e- (1)
Cathodic reaction:
2H2O + 2e- → 2OH- + H2 (2)
The general reaction of the overall corrosion process is
represented by equation (3) below.
Mg (s) + 2H2O (l) → Mg(OH)2 (s) + H2 (g) (3)
However, when Mg is exposed to chloride ions present in
the physiological environment, the Mg(OH)2 interfacial layer
reacts with the chloride ions to form highly soluble MgCl2.
The high solubility of MgCl2 and the significant reduction in
corrosion resistance provided by the reacting Mg(OH)2 layer
results in rapid dissolution of the underlying Mg substrate and
the formation of hydroxide ions and hydrogen gas [49]. The
resulting dissolution and corrosion rate are import factors in
the use of Mg as a biomaterial, since corrosion is likely to
result in mechanical failure of an implant. Therefore, the
corrosion rate must be taken into account when considering
Mg and Mg based materials for hard tissue engineering and
surgical applications. Ideally, the corrosion rate should be at a
rate that allows temporary support of tissues during the
recovery period. During the recovery period, initial
mechanical strength would be maintained until the effects of
corrosion start to occur. This would be followed by a gradual
decrease in strength over the period of tissue recovery and
finally the implant would be absorbed leaving the recovered
tissues to carry the full load [50]. Other corrosion related
factors that need to be considered is the increase in local pH
and hydrogen evolution, both of which could have significant
effects on tissues surrounding the implant. If Mg and Mg
based materials are to be successfully used as an orthopaedic
biomaterial, then the degradation behaviour and related
factors of these materials need to be effectively controlled.
3. Controlled Degradation Via Chemical
Immersion Treatment
In spite of magnesium’s many advantageous material
properties, its high chemical reactivity and poor corrosion
resistance has prevented its widespread use in orthopaedic
applications. In general, materials used in orthopaedic
application such as titanium alloys will only experience load
induced stresses in the inner core of the implant, while its
surface will be exposed and interact with the surrounding
physiological environment. Because the interfacial properties
between the implant surface and the physiological
environment are very important, different processing
techniques such as alloying, thermal spray coating, ion
implantation, micro-arc oxidation, anodizing and surface
coating treatments have been widely used to improve the
biocompatibility of the underlying material [43, 51, 52]. In the
10 Sridevi Brundavanam et al.: Chemical Immersion Coatings to Improve Biological Degradability of Magnesium
Substrates for Potential Orthopaedic Applications
case of Mg, the processing techniques have primarily focused
on improving the corrosion resistance of the metal when
exposed to the biological environment [53, 54].
Bioactive coatings such as calcium phosphate materials
have been successfully applied to a variety of metallic
implants in order to improve biocompatibility, promote
attachment to surrounding hard tissues and to suppress the
release of corrosion products into the human body [31].
Calcium phosphate (CaP) coating can be applied to a variety
of substrate materials of varying sizes and shapes using a
relatively straight forward technique known as chemical
immersion [55, 56]. Apart from convincing experimental
results of a number of independent studies that indicate CaP
coatings can significantly improve corrosion resistance
[57-59], the coatings also have the advantages of being
non-toxic, display good biocompatibility and have enhanced
bioactivity properties with respect to bone cells and other
body tissues [60]. Despite these many advantages, a number
of studies have also shown a number of shortcomings such as
poor coating adherence, surface cracking and effective control
of the CaP phases formed during immersion [54]. Regardless
of these shortcomings, biologically mimicking CaP coating
formed via the chemical immersion technique have the
potential to control the corrosion rate and enhance the
biocompatibility of Mg and Mg based alloys for orthopaedic
applications.
Figure 2. Magnesium substrate coated with a DCPD coating: (a) Optical
microscopy image of surface coating and (b) Higher resolution field emission
scanning electron microscopy micrograph showing plate-like surface
structures present in the coating.
The results of a number of recent studies have confirmed
that CaP coatings were able to enhance the corrosion
resistance and improve the biocompatibility of Mg and Mg
based alloys [61-63]. For example, a study by Xu et al. using a
chemical immersion based technique was able to form a
brushite (CaHPO4.2H2O) surface coating on an Mg-Mn-Zn
alloy substrate using an alkaline electrolyte. Regrettably,
subsequent testing revealed that the coating was porous and
failed to prevent corrosion of the substrate in a simulated body
fluid. Nonetheless, the coating did significantly reduce the
corrosion rate and provide some degree of protection against
the simulated body fluid [64]. The study not only found
improved corrosion resistance, but also found that the brushite
layer improved biocompatibility, promoted bone formation
and subsequently transformed into hydroxyapatite [Ca10(OH)2
(PO4)6] (HAP) with time [65]. In another study, Wang et al.
was able to produce a calcium-phosphate coating by
immersing an Mg substrate into a solution containing Ca and P
[Ca (NO3)2 and Na2HPO4] to form a di-calcium phosphate
di-hydrate (DCPD) surface coating [56]. The surface coating
of DCPD was effective in improving the corrosion resistance
of the substrate for the first 21 days of immersion in a
simulated body fluid. Figure 2 and the enlarged micrograph
shown in Figure 3 reveal structures in a typical DCPD coating
formed by a chemical immersion technique developed by the
authors. In a novel chemical immersion technique, Yanovska
et al. incorporated low magnetic fields into a one-step
immersion method. During this process the Mg substrate was
immersed into an aqueous electrolyte containing
Ca(NO3)2.4H2O and Na2HPO4.12H2O. The influence of
magnetic fields in orientating the crystals during the coating
formation was investigated, along with the types of phases
present. The technique produced coatings composed of DCPD
and HAP phases, the coatings were found to enhance
corrosion resistance and significantly reduced the degradation
rate of the substrate [66]. For example, Figure 4 presents a set
of representative potential-dynamic polarization curves
showing the improvement in corrosion resistance of an Mg
substrate after receiving a DCPD coating. The corrosion
testing was carried out in a phosphate buffer saline (PBS)
solution at 37 ºC as a first step in assessing the coating
bioactivity.
Figure 3. Enlarged field emission scanning electron microscopy micrograph
showing the DCPD flower-like plate structures present in the coating.
The use of HAP coatings has a number of advantages
besides improving the corrosion resistance of Mg and Mg
International Journal of Biomedical Materials Research 2014; 2(2): 7-14 11
based alloys in the physiological environment. HAP is a major
inorganic component found in natural bone tissues, therefore
using HAP as a biological coating on Mg offers a number of
attractive properties such as its good biocompatibility and
bioactivity properties with respect to bone cells and other
body tissues [67]. Other desirable properties include slow
biodegradability in situ and its ability to promote
osteoconductivity and osteoinductivity, which can accelerate
the in-growth of surrounding tissues [68-70]. These properties
are of particular importance in the case of bone tissue that are
constantly being replaced and removed by bone cells such as
osteoblasts and osteoclasts, via a process known as
remodelling. Studies have also shown that HAP displays an
excellent biocompatible response to soft tissue such as skin,
muscle and gums [71]. However, due to its low mechanical
strength HAP is restricted to low load bearing clinical
applications. Typical examples include coating the surface of
conventional metallic implants to improve their
biocompatibility and bioactivity, bone augmentation, drug
delivery, and as filler material for both bone and dental
implants [72-76].
Figure 4. Polarization curves showing the improvement in corrosion
resistance of an Mg substrate after receiving a DCPD coating. The testing
solution used was a phosphate buffer saline solution at 37 ºC and a pH of 7.4
HAP coatings can be directly deposited directly onto Mg
substrates using chemical immersion techniques or by
chemically converting modifying a pre-existing calcium
phosphate coating. For example Tomozawa et al. have treated
pure magnesium substrates with Calcium-EDTA and KH2PO4
based solutions [77]. The solution concentrations were varied
from 0.01 M/L to 0.25 M/L, while the treatment temperatures
were varied from 313 K to 373 K. Optimisation of their
immersion process revealed that by adjusting the
concentration and thermal treatment time to 0.25 M/L and 2 h,
respectively, dense rod-like HAP crystals grew along the
c-axis. HAP formation produced a dense, crystalline and
uniform coating without the formation of a Mg(OH)2
intermediate layer. The HAP coating formed by this chemical
immersion technique was found to be stable and capable of
significantly improving the corrosion resistance of the Mg
substrates. Alternatively, a two-step immersion technique can
be adopted to synthesize HAP coating on an Mg substrate.
During the first step a CaP coating is deposited onto the
substrate. Then during the second step the process converts the
calcium phosphate into HAP. For example, the authors have
produced a dicalcium phosphate dehydrate (brushite DCPD)
coating on Mg substrates from aqueous solutions containing
Ca(NO3)2 and KH2PO4 at 298 K. During the second step
DCPD coated substrates are immersed into a solution of
sodium hydroxide (NaOH) at 80 ºC for 2 h. At the end of this
low temperature thermal treatment the DCPD is converted to
HAP. The authors are currently carrying out degradation
studies to determine the corrosion resistance of DCPD and
DCPD converted to HAP coatings, Figure 4. Preliminary
results indicate that significant improvement in corrosion
resistance can be achieved by applying CaP coatings to Mg
substrates. The chemical immersion technique is a
straightforward and economic technique for synthesizing CaP
coatings such as HAP on magnesium substrates. The coatings
produced by chemical immersion were capable of slowing
down the degradation rate by significantly improving
corrosion resistance of the coated substrate.
4. Conclusion
In recent years, there has been significant research into
developing novel biologically absorbable materials for
orthopaedic applications. Mg has demonstrated that it has
some attractive properties and the potential to be used as a
biologically degradable implant material. Mg is an extremely
biocompatible material with mechanical properties similar to
bone tissue. However, in spite of having good
biocompatibility and bioactivity properties, Mg’s poor
corrosion resistance to the physiological environment has
prevented its successful use in orthopaedic applications.
Chemical immersion is an economic, efficient and
straightforward technique that offers a direct method of
depositing CaP coatings such as DCPD and HAP on Mg
substrates. Not only do CaP coatings have the potential to
reduce the corrosive effects of the physiological environment,
but they also offer the potential to significantly improve
biocompatibility and promote bone formation at the surface of
an Mg based implant. Coating Mg with a dense CaP layer that
significantly reduces the corrosion rate makes this an
attractive material for the manufacture of biodegradable
orthopaedic implants. While a significant amount of research
has been conducted in recent years investigating the potential
medical use of Mg and commercially available Mg alloys,
further research is needed to fully evaluate methods of
reducing the degradation rate in the physiological
environment. Chemical immersion is one method that has the
potential to deliver CaP coated Mg based materials with the
potential to be used in the manufacture of biodegradable
implants. However, further research is needed to fully
optimize the operational parameters of the chemical
immersion technique so that there is greater control of the
resulting coating properties. In addition, further in vitro and in
vivo studies are needed to verify the mechanical integrity of
12 Sridevi Brundavanam et al.: Chemical Immersion Coatings to Improve Biological Degradability of Magnesium
Substrates for Potential Orthopaedic Applications
the coatings and their biological compatibility with bone
tissues.
Acknowledgements
The authors would like to thank Mr Gordon Thomson for
his assistance with the optical microscopy work and Associate
Professor Gamini Senanayake for his assistance with the
electrochemical experiments in this work.
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13
2.3. Chapter summary
Chapter two was composed of two peer-reviewed papers discussing the biomedical
potential of using calcium phosphate coated Mg substrates as a potential scaffold
material in hard tissue engineering applications. The first paper consisted of a major
literature review of Mg as a potential biomaterial capable of being used in hard tissue
applications. The high corrosion rate of Mg in the physiological environment was
discussed and various current industrial treatment processes examined. Unfortunately,
many of the current treatment processes use toxic materials as additives or form part of
the coatings and surface modifications. The potential use of calcium phosphate coatings
was discussed and found to be a safe alternative to conventional surface treatments and
alloying elements. The second review paper examined the potential use of chemical
immersion as an economic, efficient and straightforward technique for reducing the
degradation rate and extending the life of an Mg based implant in the physiological
environment.
14
Chapter 3 – Case Study 1: Synthesis of Hydroxyapatite Nanometre
Scale Powders
3.1. Overview and author contributions
Chapter three addresses the first aim of the research project; namely, chemically
synthesize hydroxyapatite (HAP) in sufficient quantities from a number of sources for
chemical adsorption studies and for the development of bone cements and coatings. A
recent study by Poinern et al. has clearly demonstrated that ceramic pellets from
compacted and sintered HAP nanometre scale ultrafine powders have the capability to
significantly enhance both bioactivity and biocompatibility for potential tissue
regeneration applications [1]. Following surgical implantation into sheep M. latissimus
dorsi, the sponge-like porous structure of the pellets promoted mixed cell colonisation
and matrix deposition. Importantly, after 12 weeks their study revealed there was
extensive cellular activity present in the pellets that supported collagenous and bony
matrix deposition. The study also revealed that the pellet’s extensive three dimensional
inter-connecting pore structure promoted connective tissue and inflammatory cells to
infiltrate. The presence of different cell types, other than bone cells also confirmed
HAP’s wider biocompatibility towards other cell types. Based on the encouraging
results of Poinern’s study and similar studies reported in the literature [2-5], the present
research focused on using HAP and similar calcium phosphate materials as the primary
coating material to protect Mg substrates.
In addition to HAP’s bioactivity and biocompatibility, its complex hexagonal structure
offers an effective high capacity absorbent matrix capable of delivering a slow and
sustained release of pharmaceutical products such as antibiotics, drugs, enzymes,
15
hormones and steroids. The use of HAP as a delivery platform for pharmaceutical
products has proven to be effective in the treatment of diseases such as osteomyelitis,
osteoporosis and osseous cancer [6-9]. HAPs absorbent matrix also means that in
orthopaedic applications, like bone, it can accumulate substances like minerals. Hence,
there was a need in this research to effectively and economically produce quantities of
high quality HAP that could be used in coating and for metal ion adsorption studies
(Case Studies 2 and 3 in Chapter 4). The first part of this chapter looks at optimising a
chemical synthesis route developed by the Murdoch Applied Nanotechnology Research
Group (MANRG) at Murdoch University. And the second part looks at a new and novel
technique for converting calcium hydroxide directly into HAP (Case Study 1).
S. Brundavanam carried out all of the experiment work and conducted advance
characterisation techniques [(XRD), (SEM), (EDS) and (FT-IR)] to determine the
various physical and chemical properties of the synthesized nanometre scale HAP
powders. S. Brundavanam was also actively involved in the subsequent data analysis.
G.E.J. Poinern acted as principal supervisor and designed the overall concept for the
research forming Chapter 3 with S. Brundavanam and D. Fawcett (MANRG - Research
Fellow) acting as technical consultant. S. Brundavanam was assisted by G.E.J. Poinern
and D. Fawcett in over-coming the various technical difficulties encountered during the
research and assisted with editorial changes recommended by reviewers for the research
article presented as Case Study 1. All authors provided feedback during the preparation
of the paper.
16
3.2. Optimising Laboratory Procedure for Synthesizing Hydroxyapatite via a
combined Ultrasound and Microwave based technique.
In the present research, the MANRG developed HAP synthesis procedure was
optimised to improve its overall efficiency and product outcome. Exhaustive
descriptions of the various aspects of the MANRG procedure can be found in the
literature [1, 10, 11]. However, in the interests of completeness a brief description of the
procedure follows. Synthesis begins by adding a 40 mL solution of 0.32 M calcium
nitrate tetra-hydrate into a small glass beaker. Then approximately 2.5 mL of
ammonium hydroxide is added and mixed until a solution pH of 9.0 is achieved. The
solution is then subjected to ultrasonic irradiation for a 1 h using a Hielscher UP50H 50
W Ultrasound Processor. During the second hour of ultrasonic irradiation, a 60 mL
solution of 0.19 M potassium di-hydrogen phosphate was slowly added drop-wise.
During the second hour, the solution pH and Calcium/Phosphate [Ca/P] ratio was
maintained at 9 and 1.67 respectively. After ultrasonic treatment the solution underwent
centrifugation (15,000 g) for 20 minutes at room temperature (24 oC). The resultant
white precipitate was collected, washed and then subject to a further centrifugation for
10 minutes. The precipitate was collected and deposited into a fused silica crucible. The
loaded crucible was then thermally treated at 100 % power for 40 minutes in a domestic
microwave (1100W at 2450 MHz-LG® Australia). After thermal treatment, the dry
agglomerated powder was ball milled to produce an ultrafine nanometre scale HAP
powder. The synthesis process is presented schematically in Figure 1 to given an overall
appreciation of the procedure.
17
Figure 1. Schematic presentation of the synthesis process used to manufacture
nanometre scale hydroxyapatite powders.
3.3. Case Study 3: An alternative technique for synthesising hydroxyapatite
directly from calcium hydroxide.
Because of the advantageous biocompatible properties of HAP there is a definite need
to produce large quantities economically. Case Study 1 investigated the effects of
ultrasonic irradiation on the size, morphology and crystal structure of HAP synthesised
directly from calcium hydroxide and phosphoric acid. The study details the
experimental procedure for synthesising the nanometre scale powders and then uses
advanced characterisation techniques such as X-ray diffraction (XRD), scanning
18
electron microscopy (SEM) and Fourier transform infrared spectroscopy (FT-IR) to
investigate crystallite size, morphology and powder structure.
Published Research Article
Sridevi Brundavanam, Gérrard Eddy Jai Poinern, Derek Fawcett, Synthesis of a
Hydroxyapatite nanopowder via ultrasound irradiation from calcium hydroxide
powders for potential biomedical applications. Nanoscience and Nanoengineering.
2015; 3(1): 1-7.
Nanoscience and Nanoengineering 3(1): 1-7, 2015 http://www.hrpub.org DOI: 10.13189/nn.2015.030101
Synthesis of a Hydroxyapatite Nanopowder via Ultrasound Irradiation from Calcium Hydroxide Powders
for Potential Biomedical Applications
Sridevi Brundavanam, Gérrard Eddy Jai Poinern* , Derek Fawcett
Murdoch Applied Nanotechnology Research Group, Department of Physics, Energy Studies and Nanotechnology, School of Engineering and Energy, Murdoch University, Australia
Copyright © 2015 Horizon Research Publishing All rights reserved.
Abstract Nanoscale hydroxyapatite based ceramics are a relatively new form of materials that are currently being investigated for a number of potential biomedical applications. This study reports on a straightforward wet chemical method that uses calcium hydroxide and phosphoric acid as precursors. After chemical synthesis a conventional thermal treatment was used to produce an ultrafine hydroxyapatite nanopowder. Varying ultrasonic power between zero and 400 W during the synthesis process produced crystallite sizes ranging from 15.4 nm down to 12.2 nm. The morphology of particles synthesized under the influence of ultrasonic irradiation was predominantly spherical and granular. Also present were a small number of irregular shaped plates. Energy dispersive spectroscopy revealed the samples had a Ca:P ratio of 1.66, which was very close to the ideal value of 1.67. FT-IR studies identified functional groups and confirmed the results of the X-ray diffraction data that the powders were indeed composed of nanoscale hydroxyapatite.
Keywords Hydroxyapatite, Biomaterials, Nanometer Scale Materials, Ultrasound
1. Introduction The efficient repair of hard tissues that form the human
skeletal system continues to be a challenging goal in biomedical engineering. For many years, a variety of biologically compatible ceramics such as alumina, bioactive glasses, calcium phosphates and zirconia have been used in reconstructive and regenerative hard tissue procedures with varying degrees of success [1-3]. Importantly, any ceramic being considered for a biomedical application must be completely biologically compatible. Specifically, it must be biologically stable [4], nontoxic and non-immunogenic to body tissues [5, 6]. Among the biologically compatible ceramics, calcium phosphates are compositionally similar to
the mineral phase found in bone and consequently have been used in a variety of bone replacement therapies [7, 8]. Besides being compositionally similar, calcium phosphates are lightweight, chemical stable and are composed of ions commonly found in physiological environment [9]. The two most studied and used forms of calcium phosphates in the biomedical field are tri-calcium phosphate [Ca3(PO4)2, TCP] and hydroxyapatite [Ca10(OH)2 (PO4)6, HAP]. For example, HAP is ideal for many biomedical applications since it is the most thermodynamically stable calcium phosphate in the physiological environment [10]. Furthermore, studies have shown that HAP has good biocompatibility and bioactive properties with body tissues. In particular, for bone tissue engineering applications it offers good osteoconductivity and osteoinductivity capabilities. While it’s slow biodegradability in situ allows tissue regeneration and tissue replacement to take place [3, 11, 12]. These properties are of particular importance since bone continually undergoes cellular remodeling and as a result tissue is simultaneously deposited by osteoblasts cells and removed by osteoclasts cells [3, 13]. It is due to the advantageous bioactive properties of HAP that has made it a successful biomaterial for a number of clinical applications such as dental repair procedures, repairing bone defects, bone augmentation and coatings on metallic implants [14, 15]. The positive results achieved in these clinical applications has encouraged further research into HAP and its potential use in a variety of new tissue regeneration and tissue engineering applications [16, 17]. To date several ceramic processing techniques have been used to synthesize HAP such as wet precipitation, sol-gel, hydrothermal and ultrasonic [18-22]. However, there is a current need to develop more efficient processing techniques that have the potential to deliver large quantities of HAP nanopowders for future biomedical applications.
In the present study, a straightforward synthesis process was used to precipitate HAP seeds from a solution containing calcium hydroxide and phosphoric acid. The process was conducted at room temperature, while the solution pH was maintained at 10 via the addition of
2 Synthesis of a Hydroxyapatite Nanopowder via Ultrasound Irradiation from Calcium Hydroxide Powders for Potential Biomedical Applications
ammonia. Synthesis was carried out with and without ultrasonic irradiation to examine its influence on particle size and morphology. After a straightforward thermal treatment, particle size, composition, crystalline structure and morphology of nanopowders were studied using advanced characterization techniques such as powder X-ray diffraction (XRD) spectroscopy, Field Emission Scanning Electron Microscopy (FESEM), Energy Dispersive Spectroscopy (EDS) and Fourier Transform Infrared spectroscopy (FT-IR).
2. Materials and Methods
2.1. Materials
HAP powders were synthesised from analytical grade calcium hydroxide [Ca(OH)2] supplied the source of Ca2+ ions and analytical grade phosphoric acid [H3PO4 (86.2%)] supplied the source of phosphate ions. Calcium hydroxide was purchased from Ajax Finechem (New South. Wales, Australia) and phosphoric acid was purchased from Fisher Scientific (Fair Lawn, NJ, USA). Solution pH was achieved by the addition of ammonia that was supplied by CHEM-SUPPLY (New South Wales, Australia). All aqueous solutions were made using Milli-Q® water (18.3 MΩ cm-1) produced by an ultrapure water system (Barnstead Ultrapure Water System D11931; Thermo Scientific, Dubuque, IA). The ultrasound processor used
during ultrasonic assisted synthesis of HAP powders was an UP400S [400 W, 24 kHz, H22 Sonotrode (22 mm diameter, 45 mm maximum submerged depth)] supplied by Hielscher Ultrasound Technology. Ultrasonic outputs were controlled by operating the rotary regulator as per the manufacturer’s specifications (20% to 100%).
2.2. Synthesis of Nanometre Scale HAP Powders
The synthesis procedure begins by preparing a 100 mL aqueous solution of 0.32M Ca(OH)2 in a glass beaker. The solution was then heated to 40 ºC under vigorous stirring for 15 minutes on a standard laboratory hotplate/magnetic stirrer. Then a 100 mL solution of 0.19 M H3PO4 was slowly added to the Ca(OH)2 solution. The mixture was further stirred for 15 minutes while the mixture pH was adjusted to 10 by the drop-wise addition of ammonia. At the end of the mixing period, excess fluid was decanted from the beaker and the precipitate was deposited onto a glass dish. The dish was then placed onto a hotplate which was progressively heat up to 300 ºC. The precipitate was then thermally treated for 3 h. The synthesis procedure is schematically presented in Figure 1 (a). Ultrasonically assisted synthesis follows a similar procedure except for two 30 minutes ultrasonic treatment periods as indicated in the schematic procedure presented in Figure 1 (b). At the end of each synthesis procedure the resulting powders were examined using advanced characterisation techniques.
Figure 1. Schematic of experimental procedures used for synthesizing HAP nanopowders: (a) without ultrasound and (b) with ultrasound
Nanoscience and Nanoengineering 3(1): 1-7, 2015 3
2.3. Advanced Characterisation
2.3.1. Powder X-ray Diffraction (XRD) Spectroscopy X-ray diffraction (XRD) spectroscopy was used to study
the phase’s present and crystalline sizes in the powder samples. Spectroscopy data was recorded at room temperature, using a GBC® eMMA X-ray Powder Diffractometer [Cu Kα = 1.5406 Å radiation source] operating at 35 kV and 28 mA. The diffraction patterns were collected over a 2θ range of 20° to 60° with an incremental step size of 0.02° using flat plane geometry with 2 second acquisition time for each scan.
2.3.2. Transmission Electron Microscopy (TEM) The size and morphology of particles formed by the
synthesis process were investigated using TEM. Sample preparation consisted of collecting a portion of the synthesized powder and then placing it into a small tube containing Milli-Q® water. The tubes were then sealed and placed into an ultrasonic bath for 10 minutes. Then a single drop from the suspension was deposited onto a carbon-coated copper TEM grid using a micropipette and then allowed to slowly dry over a 24-hour period. After sample preparation a bright field TEM study was carried out using a Phillips CM-100 electron microscope (Phillips Corporation Eindhoven, The Netherlands) operating at 80kV. TEM was also used in conjunction with FESEM images to measure particle sizes and determine particle morphology
present in the powders samples.
2.3.3. Fourier Transform Infrared Spectroscopy (FT-IR)
Fourier transform infrared spectra (FT-IR) spectroscopy was used to identify species, functional groups and vibration modes present in the respective powder samples. Powder analysis was carried out using a Perkin–Elmer Frontier FT-IR spectrometer with Universal Single bounce Diamond ATR attachment. FT-IR spectra were recorded in the range from 525 to 4000 cm−1 in steps of 4 cm-1.
2.3.4. Field Emission Scanning Electron Microscopy (FESEM) and Energy Dispersive Spectroscopy (EDS)
All micrographs were taken using a FE-SEM Zeiss Neon 40 EsB with an attached energy dispersive X-ray spectrometer. The micrographs were used to study the size, shape and morphological features of the synthesized powders. Samples were mounted on individual substrate holders using carbon adhesive tape before being sputter coated with a 2 nm layer of platinum to prevent charge build up using a Cressington 208HR High Resolution Sputter coater. The EDS technique was used to verify the results of the XRD analysis and to calculate the Ca/P ratio of the synthesized powders. . In addition, FESEM was also used in conjunction with TEM images to measure particle sizes and determine particle morphology present in the powders samples.
Figure 2. XRD patterns of nanometre scale hydroxyapatite powders synthesized under the influence of ultrasonic irradiation ranging from 0 to 400 W.
4 Synthesis of a Hydroxyapatite Nanopowder via Ultrasound Irradiation from Calcium Hydroxide Powders for Potential Biomedical Applications
3. Results and Discussions
3.1. XRD Spectroscopy and TEM Analysis
Analysis of powder XRD patterns was used to identify the purity and crystalline size of HAP nanopowders produced with and without ultrasonic power. Figure 2 presents representative XRD patterns for nanopowder samples. Maximum ultrasonic power used was 400 W (100 %) and lower power level measurements were carried out at pre-determined power levels (0 %, 20 %, 40 %, 60 %, 80 % and 100 %). All patterns in Figure 2 match the known phases for pure HAP and are consistent with the phases listed in the ICDD database. In addition, each pattern identifies the main (h k l) indices associated with peaks found in the HAP samples, namely (002), (211), (300), (202), (310), (222) and (213). All patterns show characteristic HAP peaks and no evidence of non-HAP related phases were seen. Further analysis was carried out to determine the crystalline size, t (hkl), of each powder sample using the Debye-Scherrer equation [23-25].
(1)
where, λ is the wavelength of the monochromatic X-ray beam, B is the Full-Width at Half Maximum (FWHM) of the peak at the maximum intensity, θ(hkl) is the peak diffraction angle that satisfies Bragg’s law for the (h k l) plane and t(hkl) is the crystallite size. The crystallite size for each sample was calculated from the (002) reflection peak. Analysis of crystallite size data revealed increasing ultrasonic power tended to produce smaller crystallite sizes. Synthesis of powders without ultrasound power gave a mean value of 15.4 nm and powders synthesized at 400 W produced crystallites with a mean value of 12.2 nm as seen in Figure 4 (a).
Figure 3 (a) presents a representative TEM image of Ca(OH)2 particles prior to synthesis. The particles are
spherical and range in size from 0.1 µm up to 0.3 µm. Figure 3 (b) presents a typical image of HAP powders produced without the presence of ultrasonic power. The particles are rod-like to needle-like in shape and indicate growth in the c direction. The particles tend to be around 10 nm in diameter, but range in length from 50 nm upwards. However, this is not the case for powders synthesized under the influence of ultrasonic irradiation as seen in Figure 4 (b). The TEM image shows a completely different particle size and morphology. The morphology is predominantly spherical and granular in nature, but present were a small number of irregular shaped platelets which can also be seen in the high resolution FESEM image presented in Figure 4 (c).
3.2. FESEM, EDS and FTIR Spectroscopy
FESEM in conjunction with TEM microscopy was used to study particle size and morphology of nanopowders synthesized under the influence ultrasonic irradiation. Both Figure 4 (b) and (c) reveal the powder sample synthesized at 400 W is highly agglomerated. The particle morphology seen in this study is similar to morphologies previously reported in the literature [22-25]. Particle size analysis based on TEM and FESEM images for the 400 W power setting reveals a narrow particle size distribution. Particle sizes range from 20.0 nm up 66.7 nm, with spherical particles having a mean diameter of 22.8 nm and granular particles having a mean diameter of 36.7 nm as seen in Table 1. Close examination of Figure 4 (c) also reveals the presence of a small number of tube-like and irregular shaped platelets. Both of these particle features are within the narrow particle distribution as indicated in Table 1. This narrow particle distribution has also been seen in similar studies involving the sonochemical synthesis of HAP [26]. Figure 4 (d) presents a EDS spectrum of a typical powder sample showing peaks corresponding to Ca, P, and O. The presence of these peaks confirms the chemical composition of HAP. Analysis of EDS data indicated that the powder had a Ca:P ratio of 1.66, which was very close to the ideal value of 1.67.
Figure 3. (a) TEM image of Ca(OH)2 particles prior to synthesis and (b) representative TEM image of HAP synthesized without ultrasonic irradiation.
)()( θcos
9.0
hklhkl B
t λ=
Nanoscience and Nanoengineering 3(1): 1-7, 2015 5
Figure 4. (a) Graphical results of XRD analysis showing the crystallite size dependence on ultrasonic power used during synthesis, Representative TEM micrograph (b) and high resolution FESEM (c) showing the spherical and granular shaped HAP particles synthesized.
Table 1. Particle size distribution and morphology from TEM and FESEM analysis for a HAP nanopowder synthesized under the influence of 400 W.
Shape Profile Size Range (nm) Aspect Ratio Mean (nm) Std. Dev. (nm) % of Sample
Spherical (smooth)
Dia.
20.0 – 33.4 NA 22.8 3.7 44.9
Spherical (angular)
Dia.
26.7 – 46.7 NA 36.7 7.5 26.1
Irregular Width Length
20.0 – 40.2 33.4 – 66.7
1:7 27.4 47.3
7.4 4.8
17.4
Rod like Dia.
Length 20.0 – 33.3
40.0 – 66.7 1:2
21.7 43.4
4.4 8.8
11.6
Total 100.0
Figure 5 presents representative FT-IR spectra of three
HAP nanopowders synthesized under various levels of ultrasonic irradiation. The first spectrum (Blue) is a powder synthesized without ultrasonic irradiation. The other two spectra are powders synthesized under different ultrasonic power levels [Red (80 W ~ 20 %) and Green (400 W ~ 100 %). All three spectra have similar band locations and intensities normally associated with nanoscale HAP. Considering all three spectra together, starting from the right hand side of Figure 5 and moving to the left the various bands are identified. The first two bands encountered are 562 cm-1 and 601 cm-1, which are the result of ν4 vibrations produced by the O-P-O mode. The next significant bands located at 631 cm-1 and 880 cm-1 are associated with the carbonate groups and clearly indicate the presence of carbonates in all three sample spectra. The next band located at 964 cm-1 is the result of ν1 symmetric stretching vibrations
normally associated with the P-O mode. The very strong bands located at 1024 cm-1 and 1090 cm-1
indicated the presence of PO43- functional groups (P-O mode)
and the weaker bands at 1415 cm-1 and 1460 cm-1 corresponds to CO3
2- functional groups. The next band located at 1654 cm-1 also corresponds to a CO3
2- group. The presence of carbonates in the samples most likely results from atmospheric carbon dioxide interacting with the alkaline HAP precursor solution during the synthesis process [27-29]. The band located at 2978 cm-1 indicates the stretching vibrations associated with the C-H mode, while the band located at 3376 cm-1 indicates the presence of absorbed water. The final peak located at 3569 cm-1 corresponds to OH- vibrations in the HAP lattice. The results of the FT-IR analysis have clearly identified functional groups normally associated with HAP and confirm the results of the XRD data.
6 Synthesis of a Hydroxyapatite Nanopowder via Ultrasound Irradiation from Calcium Hydroxide Powders for Potential Biomedical Applications
Figure 5. FT-IR spectroscopy analysis of representative synthesized hydroxyapatite powders produced under the influence of ultrasonic irradiation.
The results of the study have shown that synthesis of nanoscale HAP from an aqueous solution containing calcium hydroxide and phosphoric acid without the aid of ultrasonic irradiation produced particles with a needle-like morphology. However, aqueous solutions subjected to ultrasonic irradiation during synthesis tended to produce a different particle size and morphology. Ultrasound assisted synthesis produced particles with a morphology that was predominantly spherical and granular in nature. Also present in the samples were a small number of irregular shaped platelets. The particle size distribution was narrow and ranged from 20.0 nm up 66.7 nm. Both XRD and FT-IR characterization studies confirmed the powders were nanoscale HAP. Further studies are needed to examine the influence of higher ultrasonic powers above the 400 W power level used in this method. The advantage of this method is its simple and straightforward set-up, and its potential scale-up capability for the production of larger powder quantities. The next stage in the research is to investigate the biological compatibility of the HAP produced using this method via in vitro studies using bone cells.
4. Conclusions The results of the study have shown that a straightforward
wet chemical method using calcium hydroxide and phosphoric acid as precursors was capable of producing HAP nanopowders. The use of ultrasonic irradiation during synthesis leads to the formation of nanoscale HAP powders with a relatively narrow size distribution. Particles produced were predominantly spherical and granular ranging in size from 20.0 nm up to 66.7 nm. There were also smaller numbers of small platelets seen in the FESEM and TEM
images. XRD studies revealed powders synthesized without ultrasonic irradiation produced a mean crystallite size of around 15.4 nm. Varying ultrasonic power from zero to 400 W produced crystallites ranging in size from 15.4 nm down to 12.2 nm. Analysis of the EDS data established a Ca:P ratio of 1.66, which was very close to the ideal value of 1.67. FT-IR studies were used to categorize functional groups present in the samples and confirmed XRD studies identifying the samples as nanoscale HAP. The study has shown that ultrasonic irradiation played an important role in determining crystallite size, particle size and particle morphology. However, further studies are needed to examine the influence of larger ultrasonic powers using this method and biological compatibility studies via in vitro cell studies needs to be undertaken.
Acknowledgements Mrs. Sridevi Brundavanam would like to acknowledge
Murdoch University for providing her PhD Scholarship to undertake the hydroxyapatite synthesis studies as part of her PhD project.
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19
3.4. Chapter summary
This chapter examined optimising the MANRG procedure for producing HAP and
developing an alternative route for producing HAP. The alternative procedure
manufactured HAP from calcium hydroxide precursor solid powders and phosphoric
acid. The procedure examined the effect of ultrasonic irradiation during the synthesis
and revealed that increasing ultrasonic power produced smaller particle sizes. The
morphology of the particles produced was spherical, granular and agglomerated. The
smallest mean particle size produced was 12.2 nm. Furthermore, the straightforward
synthesis technique has scale up potential that makes it an attractive method for
producing large quantities of HAP for commercial applications.
References
1. G. E. J. Poinern, R. K. Brundavanam, X. Thi Le, P. K. Nicholls, M. A. Cake, D.
Fawcett, The characterisation, mechanical properties and in vivo study of a
porous ceramic derived from a 30 nm sized particle based powder of
hydroxyapatite for potential hard tissue engineering applications, Sci. Rep. 4
6235 (2014) 1-9.
2. F. Despang, A. Bernhardt, A. Lode, R. Dittrich, T. Hanke, S. Shenoy, S. Mani,
A. John, M. Gelinsky, Synthesis and physicochemical, in vitro and in vivo
evaluation of an anisotropic, nanocrystalline hydroxyapatite bisque scaffold with
parallel-aligned pores mimicking the microstructure of cortical bone, J. Tissue.
Eng. Regen. Med. DOI:10.1002/term.1729 (2013).
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20
4. X. Luo, D. Barbieri, N. Davison, Y. Yan, J. D. de Bruijn, H. Yuan, Zinc in
calcium phosphate mediates bone induction: in vitro and in vivo model, Acta
Biomater. 10 (2014) 477-485.
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10. G.E.J Poinern, R. Brundavanam, X.T. Le, S. Djordjevic, M. Prokic, D. Fawcett,
Thermal and ultrasonic influence in the formation of nanometre scale
hydroxyapatite bio-ceramic, Int. J. Nanomedicine. 6 (2011) 2083-2095.
11. G. E. J. Poinern, R. K. Brundavanam, X. T. Le, D. Fawcett, The mechanical
properties of a porous ceramic derived from a 30 nm sized particles based
powder of hydroxyapatite for potential hard tissue engineering applications, Am.
J. Biomed. Eng. 2 (2012) 278-286.
21
Chapter 4 - Case Studies 2 and 3: Adsorption of Selected Heavy metals
using Hydroxyapatite absorbers
4.1. Overview and author contributions
This chapter addresses aim 2 of the research project, namely, investigating the chemical
adsorption capabilities of hydroxyapatite (HAP) to determine its suitability as an
effective bone substitute during a tissue regeneration or tissue engineering procedures
(Case Studies 2 & 3). This is of particular importance since the coating must be able to
function like the mineral phase found in bone tissue and have the ability to act as an
adsorption matrix.
S. Brundavanam carried out all of adsorption studies and conducted the advanced
characterisation techniques [(XRD), (SEM), (EDS) and (FT-IR)] needed to fully
investigate the properties of the HAP absorbers. S. Brundavanam was also actively
involved in the subsequent adsorption data analysis and interpretation of the
characterisation results. G.E.J. Poinern acted as principal supervisor and designed the
overall concept for the research forming Chapter 4 with S. Brundavanam and D.
Fawcett (MANRG - Research Fellow) acting as technical adviser. S. Brundavanam
significantly contributed to resolving various technical difficulties encountered during
the adsorption studies and assisted with editorial changes recommended by reviewers
for the two peer-reviewed research articles as presented in Case Studies 2 and 3. All
authors provided feedback during the preparation of the two articles.
22
4.2. Published Research Articles
Case Study 2
Gérrard Eddy Jai Poinern, Sridevi Brundavanam, Derek Fawcett. Kinetic and
adsorption behaviour of aqueous cadmium using a 30 nanometre scale
hydroxyapatite powder synthesized via a combined ultrasound and microwave
based technique (Currently under review - submitted to Journal of Materials and
Environmental Science)
Case Study 3
Sridevi Brundavanam, Gérrard Eddy Jai Poinern, Derek Fawcett. Kinetic and
adsorption behaviour of aqueous Fe (II), Cu (II) and Zn (II) using a 30 nanometre
scale hydroxyapatite powder synthesized via a combined ultrasound and
microwave based technique. American Journal of Materials Science 2015; 5(2):
31-40.
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Kinetic and adsorption behaviour of aqueous cadmium using a 30
nm hydroxyapatite based powder synthesized via a combined
ultrasound and microwave based technique
Gérrard Eddy Jai Poinern*1, Sridevi Brundavanam
1, Suraj Kumar
Tripathy2, Mrutyunjay Suar
3, Derek Fawcett
1
1 Murdoch Applied Nanotechnology Research Group, Department of Physics, Energy Studies and
Nanotechnology School of Engineering and Energy, Murdoch University, Murdoch, Western Australia
6150, Australia.
2 Chemical & Bioprocess Engineering Lab Center of Industrial Technology & School of
Biotechnology KIIT University, Campus-11, Bhubaneswar 751024, Odisha, India.
3 School of Biotechnology, KIIT University, Bhubaneswar-751024, Odisha, India
Received 30 January 2015
*Corresponding Author. E-mail: [email protected] Tel: (+61 8 9360-2892)
Abstract
The removal of heavy metals such as cadmium from contaminated waterways and soils is a very
important aspect of environmental remediation. This study investigated the kinetic and
absorption performance of a nanometre scale hydroxyapatite (HAP) absorber synthesised from a
combined ultrasound and microwave based technique for the removal of cadmium from an
aqueous salt solution. Parameters such as contact time, initial pH, initial cadmium concentration
and temperature were investigated. The Freundlich isotherm resulted in a more precise
modelling of the communicated experimental data. Maximum monolayer adsorption capacity of
absorber was found to be 123.45 mg/g at 298 K. Kinetic studies established cadmium
adsorption tended to follow a pseudo-second order model and intra-particle diffusion played a
significant role in determining the rate. Adsorption was endothermic, spontaneous and resulted
in structural changes to the absorber. HAP was found to be an effective absorbent material for
the removal of cadmium loaded aqueous solutions.
Keywords: nanohydroxyapatite, adsorption, cadmium, ultrasound, microwaves
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Introduction
Accumulation of heavy metal contaminants in freshwater supplies, effluents, wastewater, and
soils is an extremely important environmental problem that threatens the world. Even in small
concentrations, heavy metals are extremely hazardous to living organisms due to their toxicity
and tendency to accumulate in the food chain [1, 2]. Therefore, the removal of heavy metals
from contaminated water sources is a priority issue in many environmental remediation
programs worldwide. Heavy metals such as cadmium (Cd), copper (Cu), chromium (Cr), lead
(Pb), mercury (Hg), nickel (Ni), and zinc (Zn) are generally considered the major contaminants
of surface water, ground water and soils. The main sources of these contaminants are metal
plating industries, mining industries and drainage from abandoned disposal sites. In particular,
cadmium is highly toxic, very hazardous for aquatic and soil life and is a known carcinogen in
humans [3]. Although cadmium can be found in very small quantities in the natural
environment, a significant level of this particular harmful metal is potentially released during
industrial and mining processes [4]. Three most common modes of non-ecologically friendly
metal uptake by the human metabolic system are via nutritional intake, drinking water and
inhalation. Inhalation of cadmium produces significant irritation to the respiratory tract and all
three modes of consumption lead to anaemia, osteoporosis, osteomalacia, kidney damage and
Itai-itai disease [5-7]. In the case of osseous related diseases, it is the remarkable ion exchange
capacity of natural hydroxyapatite (HAP) present in bone tissue that permits the replacement of
calcium ions with cadmium ions.
It is due to this extraordinary ability of HAP to accumulate and bind with heavy metals
ions such as cadmium within the body, which produces the serious detrimental effect on the
health and well-being of an individual. In the case of environmental remediation of
contaminated wastewaters, several processes such as filtration, chemical precipitation,
electrochemical deposition, ion exchange, adsorption and solvent extraction have all been
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extensively used with varying degrees of efficiency and cost effectiveness [8-10]. Chemical
immobilization by adsorption is one method for reducing the bioavailability of toxic metals via
the formation of new stable minerals with lower solubilities in the environment. Several studies
have identified the high adsorption capacity of HAP and its ability to combine especially with
divalent heavy metals. It is this property that enables HAP to significantly reduce metal ion
concentrations in aqueous solutions [11-14]. HAP is a naturally occurring mineral with a
hexagonal crystal structure composed of calcium phosphate groups and has the general formula
of [Ca10(OH)2 (PO4)6] for the unit cell. Both chemical and crystallographic studies have also
shown that synthetic HAP is similar to the chemical composition of the naturally occurring
inorganic component found in teeth and bone tissue [15]. Synthetic forms of HAP have been
successfully used in a wide range of applications such as bioceramics for dental and bone repair
procedures, absorbents to separate enzymes and proteins during chromatography and as
catalysts for dehydrogenation and dehydration of alcohols [16, 17]. The wide range of
applications stem from the advantageous surface properties of HAP that include a hydrophilic
nature, surface charge, pH, porous structure and 2.6 P-OH surface groups per nm2 which act as
sorption sites [18, 19]. It is the sorption properties of HAP that makes it an attractive absorbent
material for potential remediation of heavy metal contaminated water, which adds up from the
environment naturally or synthesized and disposed as waste industrial process by-products into
the ecosystem.
The synthesis of nanometre scale crystalline HAP (nano HAP) for use as an adsorbent
material has been extensively studied and has resulted in techniques such as emulsions,
hydrothermal and solution-gelation being used to produce the material [20-22]. However, wet
chemical synthesis offers many advantages due to its straightforward and economically efficient
protocols during the process. In particular, one advantage is being able to control size and
morphology of the forming particles by varying experimental parameters. It is the control of
experimental parameters that directly regulates particle nucleation, aging and growth kinetics.
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Recent studies have also investigated the use of ultrasonic irradiation during wet chemical
synthesis to enhance the manufacture of nanometre scale materials [23, 24]. During ultrasonic
irradiation, acoustic cavitations create bubbles that grow and then implosively collapse creating
localized hot spots with temperatures as great as 5300 K and pressures around 500 atmospheres.
The acoustic cavitations also produce very rapid cooling rates that often exceed 1010 K/s [25].
It is the extreme pressures and temperatures experienced during ultrasonic irradiation that
promotes the physical effects and chemical reactions which directly influence the size and
morphology of the forming particles [26]. Another technique that has been successfully used to
improve the synthesis route is microwave heating that significantly reduces reaction times and
increases product yields when compared to conventional heating methods [15, 26, 27].
Therefore, incorporating ultrasonic and microwave based techniques into the wet chemical
method offers greater efficiency in manufacturing nanometre scale materials. In this study,
solutions containing calcium, hydroxyl and phosphate ions were subjected to ultrasound
irradiation to form the initial calcium phosphate compounds. This was followed by thermally
treating the compounds in a microwave oven to produce the nanometre scale HAP powders
used in characterisation and adsorption studies.
The objectives of this study were to first synthesize a nanometre scale HAP powders via
a combined ultrasonic and microwave heating based wet chemical method. And secondly,
investigate the potential use of the powders as an adsorbent material for the removal of Cd2+
cations from aqueous solutions. The synthesized powders were characterized using X-ray
diffraction (XRD), field emission scanning electron microscopy (FESEM), energy dispersive
spectroscopy (EDS) and Fourier transform infrared spectroscopy (FTIR). The adsorption
capacity of the powders was investigated via the removal of Cd2+
cations from cadmium
contaminated water using a batch equilibrium procedure. Both Langmuir and Freundlich
adsorption isotherms were used to model the experimental data. While the kinetic behaviour of
the adsorption mechanism were studied using Lagergren’s pseudo-first order, McKay & Ho’s
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pseudo-second order and intra-particle diffusion models. Furthermore, the influence of initial
Cd2+
cation concentration, contact time, temperature, solution pH and thermodynamic
parameters were all evaluated from the adsorption measurements.
2. Materials and methods
2.1. Materials
HAP powders were synthesized from high purity calcium nitrate tetra-hydrate [Ca(NO3)2.4H2O]
and potassium di-hydrogen phosphate [KH2PO4], while solution pH was controlled by the
addition of ammonium hydroxide [NH4OH]. An Ultrasound Processor [UP50H: 50 W, 30 kHz,
MS7 Sonotrode (7mm diameter, 80 mm length)] supplied by Hielscher Ultrasound Technology
was used to deliver ultrasound irradiation during HAP synthesis. The source of Cd 2+
ions used
in the adsorption studies was high purity cadmium chloride [CdCl2]. All chemicals used in this
work were supplied by Chem-Supply (Australia) and all aqueous solutions were made using
Milli-Q® water (18.3 M cm
-1) produced by an ultrapure water system (Barnstead Ultrapure
Water System D11931; Thermo Scientific, Dubuque, IA).
2.2. Synthesis of nanometres scale hydroxyapatite
The synthesis procedure used to produce the nanometre scale HAP powders was developed by
the authors and a detailed description is given elsewhere [15, 24 &28]. However, in the interest
of completeness a brief description is presented. The procedure in brief consists of adding a 40
mL solution of 0.32 M calcium nitrate tetra-hydrate into a small glass beaker and then adjusting
the solution pH to 9.0 using approximately 2.5 mL of ammonium hydroxide. The solution was
then sonicated for 1 h using the Ultrasound Processor set to 50 W and maximum amplitude.
During the second hour of ultrasound treatment 60 mL of 0.19 M potassium di-hydrogen
phosphate solution was slowly added while the solution pH was maintained at 9.0 and the
Calcium/Phosphate [Ca/P] ratio was maintained at 1.67. On completion of the ultrasound
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treatment, the solution was then subjected to 20 minutes of centrifugation (15,000 g) at room
temperature to produce a solid white precipitate. The precipitate was collected, washed and
centrifuged for a further 10 minutes. After the second centrifugation, the precipitate was
deposited into a fused silica crucible before being placed into a domestic microwave oven for
thermal treatment [Set at 100 % power for 40 minutes: 1100W at 2450 MHz-LG® Australia].
After thermal treatment, the resultant agglomerated powder was then ball milled until an
ultrafine nanometre scale HAP powder was produced.
2.3. Advanced characterisation techniques
2.3.1. X-ray diffraction (XRD) spectroscopy
Powder X-ray diffraction (XRD) spectroscopy was used to examine and identify the purity,
crystalline size and phases present in synthesized nanometre scale HAP powders. Spectroscopy
data was recorded at room temperature, using a Siemens D500 series diffractometer [Cu Kα =
1.5406 Å radiation source] operating at 40 kV and 30 mA. The diffraction patterns were
collected over a 2θ range of 20° to 60° with an incremental step size of 0.04° using flat plane
geometry with 2 second acquisition time for each scan. The crystalline size of the particles in
the powders was calculated using the Debye-Scherrer equation [Equation 1] from the respective
spectroscopy patterns.
2.3.2. Field emission scanning electron microscopy (FESEM)
FESEM was used to study the size, shape and morphological features of the HAP powders
before and after the adsorption studies. All micrographs were taken using a high resolution
FESEM [Zeiss 1555 VP-FESEM] at 3 kV with a 30 µm aperture operating under a pressure of
1.33310-10
mbar. Samples were mounted on individual substrate holders using carbon adhesive
tape before being sputter coated with a 2 nm layer of gold to prevent charge build up using a
Cressington 208HR High Resolution Sputter coater.
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2.3.3. Energy dispersive spectroscopy (EDS)
EDS was used to provide an elemental analysis of the powder samples using an Oxford
Instruments energy dispersive system (133 eV resolution), via a 10 mm2 SATW detector. The
analysis was carried out to verify the results of the XRD analysis and to calculate the Ca/P ratio
of the synthesised nano-HAP powders.
2.3.4. Fourier transform infrared spectroscopy (FT-IR)
FT-IR spectroscopy of synthesized nanometre scale HAP powders and cadmium loaded HAP
powders was carried out using a Perkin–Elmer Frontier FT-IR spectrometer with Universal
Single bounce Diamond ATR attachment. Both FT-IR spectra were recorded in the range from
525 to 4000 cm−1
in steps of 4 cm-1
.
2.4. Batch adsorption studies
All adsorption experiments were carried out using the batch equilibrium technique. During
experiments the adsorption capacity of absorber for Cd2+
ions was investigated. The study also
examined the influence of the initial Cd2+
ion concentration, contact time, temperature and pH
of the test solutions.
The influence of contact time on Cd2+
ion adsorption on absorber was examined using
aqueous solutions containing 100 mg/L of Cd2+
ions (100 ppm) prepared from cadmium
chloride [CdCl2]. A sample of 0.1 g HAP taken from the stock solution (1g/L) was added to a
100 ppm Cd2+
prepared solution. The magnetic stirring speed was set to 400 rpm, while the
temperature of the suspension was maintained at 298 ± 1 K. The pH of the suspension was
adjusted by adding drops of 0.1 M NaOH to the suspension and then maintaining the value of 7
throughout the experiment. Sample volumes were taken from the suspension during the mixing
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process at pre-determined time intervals (10, 20, 30, 40, 60, 90, 120, 180, 240 and 300 min) so
that Cd2+
ion concentration in the solution could be measured. After each specified time, the
sample solution was first filtered using a Whatman®
0.22μm membrane syringe filter and then
centrifuged at 15,000 g for 20 minutes. Each time a sample volume was taken, an equal volume
was then added to the solution to maintain the initial volume. The concentration of Cd2+
ions in
the sample solution was determined by inductively coupled plasma atomic emission
spectrometer (Varian Vista Axial CCD ICP-AES). All experiments were carried out in
triplicate.
The effect of solution pH on adsorption was studied by adjusting the pH from 2 to 12.
This was done by treating the solution with either 0.1M HCl or 0.1M NaOH. The temperature
was normally set to 298 K, except where temperature variation studies were carried out. After
300 min, the solid was separated from the suspension using a Whatman® 0.22μm membrane
syringe filter and the residual cadmium level was measured. The influence of initial Cd2+
ion
concentration was studied by first preparing a series of Schott reagent bottles containing 100
mL aqueous solutions consisting of varying concentrations of Cd2+
ion (50 to 175 mg/L) with an
initial pH of 7. Then 1 g/L stock solution containing nanometre scale HAP powder was added
(100 mL) into each bottle. The bottles were then sealed and the magnetic stirring speed was set
to 400 rpm, while the bottles were thermostatically maintained at the respective isotherm (283,
293, 303, 313 and 323 K). Measurement of Cd2+
concentration levels was carried out at pre-
determined time intervals (10, 20, 30, 40, 60, 90, 120, 180, 240 and 300 min). All initial Cd2+
ion concentration experiments were carried out in triplicate. The data collected from the
adsorption experiments was then used in the subsequent kinetic and adsorption isotherm
modelling studies.
3. Results and discussion
3.1. XRD spectroscopy analysis
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Analysis of respective powder XRD patterns was used to identify the purity and crystalline size
of the synthesised nanometre scale HAP powders used in the adsorption studies. A XRD pattern
of a representative pure HAP powder before being used in the adsorption studies is presented in
Figure 1. Inspection of the HAP powder (blue pattern in Figure 1) reveals that the reflection
pattern matches the known phases present in pure HAP and is consistent with the phases listed
in the ICDD database. The pattern also identifies main (h k l) indices found in the HAP sample,
namely (002), (211), (300), (202), (310), (222) and (213). The pattern only shows characteristic
HAP peaks and there was no evidence of non-HAP phases. In addition, the crystalline size, t (hkl),
of each sample was calculated from the respective XRD patterns using the Debye-Scherrer
equation [29-31]
)(
)(θcos
9.0
hkl
hklB
t
(1)
where, λ is the wavelength of the monochromatic X-ray beam, B is the Full Width at Half
Maximum (FWHM) of the peak at the maximum intensity, θ(hkl) is the peak diffraction angle
that satisfies Bragg’s law for the (h k l) plane and t(hkl) is the crystallite size. The crystallite size
calculated from the (002) reflection peak for the sample gave a mean value of 30 nm.
Studies have shown that HAP has a strong selectivity towards divalent metal cations via
an ion-exchange mechanism [19, 32]. The results of this study confirm HAPs selectivity
towards Cd2+
cations and the substitution of Cd2+
cations for lattice Ca2+
cations. The ion-
exchange mechanism was also assisted by the slight difference between the respective ionic
radius’s of Cd2+
(9.7 nm) and Ca2+
(9.9 nm). After each adsorption procedure the solid residue
was collected and subjected XRD analysis. Inspection of the XRD pattern for a representative
Cd2+
loaded sample is presented in Figure 1 and reveals structural changes resulting from Cd2+
absorption. The ion-exchange that takes place between the Cd2+
and Ca2+
cations also results in
a very small decrease in the unit cell dimensions and hence volume of the unit cell. Thus, the
associated small XRD shifts reported in the literature were also seen in this study and support
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the ion-exchange mechanism [11, 33]. In this study, the ion-exchange mechanism between Cd2+
and Ca2+
can be represented by the equivalent molar exchange in general HAP formula Ca10−x
Cdx (PO4)6 (OH)2, where x can vary from 0 to 10 depending on experimental parameters and
reaction time [11].
Figure 1: XRD pattern of synthesized nanometre scale HAP powder after milling and drying
(Blue), while the second XRD pattern (Red) reflects the maximum uptake of Cd2+
of the solid
residue at the end of the adsorption study. The study used 1 g/L HAP absorbent, a 100 mg/L
Cd2+
ion solution, with a pH of 7 and stirred at 400 rpm for a total contact time of 300 min.
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3.2. FESEM and EDS analysis
FESEM was used to study particle size and morphology before and after adsorption studies. A
typical micrograph of a pure HAP powder sample is presented in Figure 2(b) and reveals a
sphere like particle morphology that is highly agglomerated. The spherical morphology seen is
similar to particle morphologies previously reported in the literature [15, 29-31]. Particle size
analysis of micrographs revealed a mean particle diameter of 28 ± 5 nm and compared
favourably to the XRD determined value of 30 ± 5 nm. After adsorption studies, FESEM
analysis was also performed on samples to determine any changes in size or morphology. Figure
2 (d) presents a typical micrograph of the Ca10−xCdx(PO4)6(OH)2 structure displaying the same
agglomerated spherical morphology seen in samples prior to adsorption studies. The mean
particle size seen in the Cd2+
loaded samples was estimated to be 26 ± 5 nm. The studies have
revealed that there was no significant difference in particle size and morphology between
unloaded and loaded powder samples. Figure 2 (a) presents a typical EDS spectrum of an
unloaded sample showing peaks corresponding to Ca, P, and O, and confirms the chemical
composition of HAP. In addition, a number of Au and Cu peaks were also seen in the EDS
spectrum. The Au peaks were the product of the Au coating used on the samples, while the Cu
peaks are a consequence of X-rays being scattered from the copper grid. Figure 2 (c) presents an
EDS spectrum of a representative Cd2+
loaded sample and confirms the chemical composition of
the sample. The presence of two cadmium peaks (indicated by red arrows) also confirms the
results of the XRD analysis discussed above.
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Figure 2: EDS spectrum (a) and FESEM micrograph (b) of unloaded nanometre scale
crystalline HAP powder and EDS spectrum (c) and FESEM micrograph (d) of a typical Cd2+
loaded HAP residue.
3.3. FT-IR studies
FT-IR spectroscopy and subsequent analysis was used to identify species, functional groups and
vibration modes associated with peaks seen in sample spectra taken before and after adsorption
studies. Figure 3 presents the results of FT-IR spectroscopy of representative powder samples
taken before and after adsorption studies. Starting from the right hand side of Figure 3 with a
powder sample prior to adsorption testing. We see three peaks occurring at 561 cm-1
, 600 and
631cm-1
that are consistent with 4 vibrations normally associated with O-P-O modes. The
weaker peak located at 832 cm-1
is associated with a carbonate group and clearly indicates the
presence of carbonates in the sample. The presence of carbonate ions in the sample is a
consequence of atmospheric carbon dioxide interacting with HAP precursors in the synthesis
solution and has been reported in the literature by other researchers [34, 35]. The small peak
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located at 963 cm-1
indicates 1 symmetric stretching vibrations normally associated with a P-O
mode. The strong peak located at 1027 cm-1
and the smaller peak located at 1091 cm-1
correspond to PO43-
functional groups. While the weaker peak located at 1374 cm-1
corresponds
to a CO32-
functional group. Moving further leftward we encounter a smaller peak located at
1644 cm-1
which indicates the presence of a CO32-
group. The next peak located at 3215 cm-1
indicates the presence of absorbed water, while the last identified weak peak located at 3570 cm-
1 corresponds to OH
- ion vibrations in the HAP crystal lattice. The second pattern presents the
FT-IR analysis of a representative powder sample after adsorption testing. The results are very
much the same as the pre-adsorption sample, except both 600 and 631cm-1
peak intensities are
significantly smaller than those of the pre-adsorption sample as seen in Figure 3 insert. Both
peaks are associated with the P–OH groups that act as sorption sites on the surface of the HAP
powder [18, 19]. Collectively, their reduced intensities confirm the decrease in the number of
free sites available due to Cd2+
ion attachment and collaborates the results of XRD analysis.
3.4. Effect of initial Cd2+
concentration and pH on adsorption
The influence of initial Cd2+
concentration and contact time on adsorption was carried out on
initial cadmium concentrations ranging from 50 to 175 mg/L and over contact times ranging
from 10 to 300 minutes. Solution temperature was maintained at 298 ± 1 K and solution pH was
maintained at 7. During adsorption studies all solutions were magnetically stirred at 400 rpm
and concentration measurements were taken at different interval times over a 300 minute test
period (10, 20, 30, 40, 60, 90, 120, 180, 240 and 300 min).
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Figure 3: FT-IR Spectroscopy analysis of synthesized pure nanometre scale hydroxyapatite
powder after drying and a cadmium loaded HAP residue. The study used 1 g/L HAP absorbent,
a 100 mg/L Cd2+
ion solution, with a pH of 7 and stirred at 400 rpm for a total contact time of
300 min.
Figure 4 (a) presents the adsorption results of a HAP absorber (1 g/L) placed in an initial Cd2+
ion concentration of 100 mg/L maintained at 298 ± 1 K and stirred at 400 rpm for 300 min.
Inspection of Figure 4 (a) reveals maximum Cd2+
removal (75 %) occurred after 90 minutes and
beyond this period no further adsorption was observed. The trend was typical of the absorber
and further investigations were undertaken to quantify Cd2+
ion uptake by increasing initial
metal concentrations. The results of these investigations are presented in Figure 4 (b) and reveal
increasing initial Cd2+
ion concentration (50 to 175 mg/L) tended to produce a corresponding
increase in capturing capacity of the absorber (45.75 to 110 mg/g). The increasing trend
suggests the higher initial concentrations were able to overcome mass transfer related
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resistances existing between the aqueous and solid absorber phase by effectively creating a
driving force.
The effect of pH on adsorption was investigated by repeating batch equilibrium studies
using an initial Cd2+
ion concentration of 100 mg/L (100 ppm) at various pH values ranging
from 2 to 12. Figure 4 (c) presents the results of the pH study and illustrates the effect of pH on
adsorption. It is apparent form Figure 4 (c) that Cd2+
adsorption steadily increases from pH 2 up
to a value of 12 and pH influences the degree of Cd2+
removal from the solution. At lower pH
values H+ ions compete with Cd
2+ ions for binding sites on the absorber surface. However, as
the pH increases there is a reduction in competition between H+, Cd
2+ and Cd(OH)
+ ions and as
a result, metal uptake by the absorber increases [11]. The experimental results indicate Cd2+
removal was predominantly controlled by adsorption up to a pH of around 8. However, beyond
pH 8 Cd2+
ion removal was significantly enhanced by cadmium hydroxide precipitation.
3.5. Adsorption kinetics
Kinetic analysis is essential for understanding any adsorption process. This is because sufficient
residence time on the absorber surface is needed to complete the adsorption reaction. The
amount of Cd2+
ions adsorbed at equilibrium time (qe) was calculated using equation 2 below:
(2)
where Co and Ce are the initial and equilibrium concentrations (mg/L) of Cd2+ ions in solution,
V is solution volume (L) and m is absorber mass (g) used during the experiments. The
experimental adsorption data was analysed using three kinetic models: 1) Lagergren’s pseudo-
first order law; 2) McKay and Ho’s pseudo-second-order law, and 3) the intra-particle diffusion
model. In the first case, the Lagergren pseudo-first order law [36] is defined by equation 3
below:
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(3)
where, qt (mg/g) is adsorption at time t and k1(min-1
) is the pseudo-first order adsorption rate
constant. In the second case, adsorption kinetics was examined using the McKay and Ho’s
pseudo-second-order law [37], which is described by equation 4 below:
(4)
where k2 (g/min.mg) is the pseudo-second-order rate constant for adsorption. Figure 5 (a)
presents kinetic data plotted using the Lagergren pseudo-first order equation (3), while Figure 5
(b) presents kinetic data plotted using the McKay and Ho’s pseudo-second-order equation (4).
Inspection of the two plots reveals that they are both linear in nature and both models are
reasonable representations of the kinetic data. In spite of the reasonable agreement, further
analysis reveals that the second order model is superior and has a slightly higher correlation
coefficient (R2) as seen in Table 1.
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Figure 4: (a) Influence of initial Cd2+
concentration and contact time; (b) capturing
capacity of the HAP absorbent with increasing initial Cd2+
ion concentration (50 to 175
mg/L); (c) the effect of pH on Cd2+
adsorption, and (d) the effect of temperature % Cd2+
removal.
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Figure 5: Adsorption data modelled using three kinetic models: (a) Lagergren’s pseudo-first
order law; (b) McKay and Ho’s pseudo-second-order law, and (c) the intra-particle diffusion
model. [The study used 1 g/L HAP absorbent, a 100 mg/L Cd2+
ion solution, with a pH of 7 and
stirred at 400 rpm for a total contact time of 300 min].
Table 1: A comparison between the pseudo kinetic (first and second order) rate constants and
intra-particle kinetic diffusion constants and calculated qe values (pH = 7, initial Cd2+
concentration = 100 mg/L, adsorbent dosage = 1 g/L, and agitation rate = 400 rpm).
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Pseudo-first-order kinetic model
Temperature (K) k1 (min-1
) qe (mg/g) R2
298 0.0317 67.17 0.9964
Pseudo-second-order kinetic model
Temperature (K) k2 (g/mg.min) qe (mg/g) R2
298 0.00076 80.64 0.9972
Intra-particle kinetic diffusion model
Temperature (K) kp (g/mg.min) C (mg/g) R2
298 7.726 4.448 0.9764
However, the adsorption process is not a single step event, but instead involves a
number of steps. The first step involves diffusion of sorbate from the aqueous solution to the
absorber surface. This is followed by the much slower diffusion of sorbate into the internal
porous voids of the absorber matrix. Weber and Morris [38] have described this multi-step
process using an intra-particle diffusion model which is described by equation (5) below:
(5)
where C is the intercept that provides the ideal boundary layer thickness and kp is the intra-
particle diffusion rate constant (mg/g.min1/2
g). In the model, the resulting plot of qt versus t1/2
will be linear if the intra-particle diffusion process is involved. In addition, if the linear line
passes through the graphs origin, it indicates there is only one rate-controlling step. On the other
hand, if the graph is composed of a series of linear plots, then each linear plot will represent a
different stage in the adsorption process. The first plot is generally steeper and results from the
rapid diffusion of sorbate from solution and its subsequent surface attachment to the absorber.
The second linear portion of the graph has a smaller, more gradual gradient and indicates intra-
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particle diffusion is the prevailing process. In the final equilibrium stage, the gradient levels off
and indicates a significant reduction in diffusion due to the low solute concentration. When the
adsorption data was plotted using the intra-particle diffusion model a typical multi-gradient line
graph pattern was present and is presented in Figure 5 (c). For clarity, only the second and final
equilibrium stages are plotted and the modelling parameters such as rate constants, intercepts
and correlation coefficients are presented in Table 1.
3.6. Adsorption isotherms
Analysis of the experimental data was important in determining the distribution of Cd2+
ions
between solution and HAP absorber when the adsorption process had reached its equilibrium
state. There are a number of isotherm equations available to model the results of equilibrium
data obtained from adsorption systems. The two most widely used equilibrium modelling
equations are Freundlich and Langmuir. Freundlich is purely an empirical equation used to fit
experimental data. It takes into account surface heterogeneity, the exponential distribution of
active adsorption sites and their respective energies over a wide range of concentrations.
Langmuir, unlike the Freundlich assumes maximum adsorption occurs when the surface is
covered by a monolayer of adsorbate. The equilibrium data for Cd2+
ions in solution over a
concentration range starting from 50 to 175 mg/L at constant temperature of 298 K, pH of 7, 1
g/L absorbent dose and a contact time of 300 minutes was analysed using Freundlich and
Langmuir isotherms. The Freundlich isotherm equation used for modelling the equilibrium data
is presented in its linear form below:
(6)
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where kF and n are Freundlich parameters related to the extent of adsorption and the intensity of
adsorption respectively. Both kF and n parameters were determined via plotting log qe versus log
Ce , the results of which are presented in Figure 6 (a).
The Langmuir isotherm used for modelling the equilibrium data is expressed
mathematically in the linear form by equation (7) below:
(7)
where, Qm (mg/g) is the monolayer adsorption capacity, b (L/g) is the Langmuir constant that is
related to the free energy of adsorption, Ce (mg/L) and qe (mg/g) are the equilibrium
concentrations of adsorbate in solution and on the surface of absorber respectively. Using the
Langmuir isotherm, a linear plot of equilibrium data was obtained when Ce/qe was plotted
against Ce over the entire Cd2+
ion concentration range as seen in Figure 6 (b). The adsorption
isotherm is characterized by the two parameters Qm and b that were determined from the plot.
The parameters reflect the surface properties and affinity of the Cd2+
ions for the HAP absorber.
Both the Freundlich and Langmuir isotherm plots displayed good linear fits and were able to
provide intercept and slope parameters which are listed in Table 2. For the Freundlich isotherm
a value of 3.106 was obtained for n, which fell within the range of 1 to 10 and indicated
adsorption was favourable. Furthermore, based on the coefficient of correlation (R2) values, the
Freundlich isotherm provided the best fit of the experimental data. However, unlike Langmuir,
the Freundlich isotherm does not predict any maximum occupancy of Cd2+
ions on the HAP
absorber surface. Instead, it mathematically predicts infinite surface occupancy and the
possibility of multi-layered adsorption. Nonetheless, the maximum monolayer adsorption
capacity (Qm) calculated from the Langmuir equation was found to be 123.45 mg/g. Analysis of
isotherm data indicates adsorption capacity was the result of increasing equilibrium cadmium
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concentrations in solution. The increased concentrations were able to increase the numbers of
Cd2+
ions at the absorber surface and enhance the probability of adsorption.
Figure 6: Linear fits of experimental data using (a) Freundlich and (b) Langmuir isotherms
[The study used 1 g/L HAP absorbent, a 100 mg/L Cd2+
ion solution, with a pH of 7 and stirred
at 400 rpm for a total contact time of 300 min].
In addition, a comparative evaluation of the Cd2+
adsorption efficiency of the nanometre
scale HAP absorber used in the present study and several other materials reported in the
literature is made via Table 3. Inspection of Table 3 reveals the HAP absorber used in this study
possesses a higher adsorption capacity to many other materials reported in the [11] listed in the
table. But only activated carbon base materials such as Filtrasorb 400 have higher adsorption
capacities [46].
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3.7. The effect of temperature
An isothermal study was also carried out to investigate the effect of temperature on cadmium
adsorption of HAP absorber. The five isotherms used were 283, 293, 303, 313 and 323 K, while
the initial Cd2+
ion concentration (100 mg/L), contact time (300 min) and adsorbent dose (1 g/L)
were kept constant during the study. The study revealed cadmium adsorption capacity steadily
increased with increasing temperature [see Figure 4 (d)] and indicated the adsorption process
was endothermic in nature.
To investigate the thermodynamics of the adsorption process, parameters such as free
energy change (ΔG0), enthalpy change (ΔH
0) and entropy change (ΔS
0) were calculated using
the following equations:
KRTG ln0 (8)
000 STHG
(9)
For the purpose of producing a linear graphical representation, equations (8) and (9) were
combined and rearranged to give Van’t Hoff equation:
RT
H
R
SK
00
ln
(10)
where R is the universal gas constant (8.314 J mol-1
K-1
), T is the temperature in Kelvin and K is
the thermodynamic equilibrium constant in units of L/g and is generally expressed by:
e
e
C
qK
(11)
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In order to ensure that K in equation (11) is dimensionless we have adopted the method
suggested by Milonjic in which the equilibrium constant K is replaced with a new dimensionless
equilibrium constant Kd [39]:
(12)
where ρ is the density of water (~1000 g/L) and is assumed to remain constant over the entire
temperature range studied and contains only a small absorber dose.
Table 2: Comparisons between Freundlich and Langmuir adsorption isotherm constants for
Cd2+
onto nanometre scale hydroxyapatite at 298 K.
Freundlich adsorption isotherm constants
Temperature (K) kf (mg/g) n R2
298 28.47 3.106 0.9967
Langmuir adsorption isotherm constants
Temperature (K) Qmax (mg/g) b (L/mg) R2
298 123.45 0.091 0.9868
Table 3: Comparison between cadmium absorption capacities of various materials
Absorbent qm (mg/g) Reference
Rice husk 2.00 [40]
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Olive cake 65.40 [41]
Sugarcane bagasse 24.70 [42]
Cashew nut shell 14.29 [43]
Coffee grounds 15.65 [44]
Calcium Alginate Beads 32.94 [45]
Activated carbon (Filtrasorb 400) 307.50 [46]
Nanometre scale HAP
(Needle shape 20 - 30 nm)
142.86 [11]
Nanometre scale HAP
(Spherical particles 30 nm)
123.45 Present Study
Values of Kd were calculated from the equilibrium concentrations over the temperature
range. The values were found to increase with increasing temperature and revealed adsorption
was endothermic in nature. The calculated changes in Gibbs free energy ∆G0 were plotted
against temperature T (K) and from the graph values of ∆H0 (slope) and ∆S
0 (intercept) were
determined (Graphs not shown). However, the thermodynamic parameters derived from this
analysis are presented in Table 4. Inspection of Table 4 reveals that each of the ΔG0 values are
negative for each isotherm. The result indicates cadmium adsorption is spontaneous in nature
and the adsorption process is enhanced with increasing temperature. The positive value of ΔH0
(6872.12 J mol-1
) indicates the endothermic nature of the adsorption process and the positive
value of ΔS0 (90.06 J mol
-1 K
-1) indicates the affinity of the nanometre scale HAP absorber for
Cd2+
ions. The positive values of both ΔH0 and ΔS
0 are also responsible for making the reaction
spontaneous, irreversible in nature and also suggest that there were some structural changes to
the absorber. Throughout Cd2+
ion removal from solution two distinct stages were clearly seen.
During the first stage P-OH surface groups located on the absorber surface promoted surface
complex’s to take place. The FT-IR spectroscopy analysis presented in Figure 3 confirms this
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stage. Since the 600 and 631cm-1
peak intensities normally associated with P–OH groups were
significantly reduced after Cd2+
ion adsorption had taken place. The reduction in peak intensity
reflects the significant reduction in the number of free sites available for attachment to take
place. During the second stage significant numbers of Cd2+
ions were substituted for Ca2+
ions in
the HAP crystal lattice. The presence of cadmium in both the XRD and EDS analysis confirms
the ion-exchange mechanism. The adsorption process can be described by the following
equation:
Ca10 (PO4)6 (OH)2 + xCd2+
→ Ca10−x Cdx (PO4)6(OH)2 + xCa2+
(13)
It is during the second stage when there is partial dissolution of calcium from the HAP lattice
that results in a apatite precipitation described by Ca10−x Cdx (PO4)6(OH)2 and Ca2+
ions.
Table 4: Thermodynamic parameters for Cd2+
adsorption onto nanometre scale hydroxyapatite.
Temperature
(K)
Kd ΔGº
(J/mol)
ΔHº
(J/mol)
ΔSº
(J/mol K)
283 2773.58 -18653.24
293 3000.00 -19503.53
303 3255.32 -20374.94 6872.12 90.06
313 3545.45 -21269.55
323 4000.00 -22273.02
J. Mater. Environ. Sci. 6 (Y) (2015) xxx-xxx
Poinern et al.
ISSN : 2028-2508
CODEN: JMESCN
49
Conclusion
The present study has shown a HAP powder synthesized via a combined ultrasound and
microwave based technique has resulted in the development of a highly crystalline structure
with a spherical mean particle size of 30 nm. The prepared nanometre scale HAP powder was
found to be an effective adsorbent for the removal of Cd2+
cations from aqueous solutions. The
sorption performance was found to be a function of initial Cd2+
concentration, temperature and
solution pH. Furthermore, Cd2+
removal was found to improve with increases in these
parameters for specific contact times. The study also confirmed the sorption process was
endothermic in nature and Cd2+
sorption increased with temperature. Kinetic studies revealed
the sorption process closely followed pseudo-second order kinetics. During sorption, the initial
uptake rate of Cd2+
was high, but this was followed by a much lower uptake rate. The ion-
exchange mechanism (Cd2+
→ Ca2+
) was clearly identified as the major participant in the
sorption process. The results of both XRD and EDS analysis confirmed ion-exchange was a
major removal mechanism and sorption was heavily influenced by intra-particle diffusion.
Isotherm studies indicated the Freundlich isotherm modelled experimental data better than the
Langmuir isotherm. However, the Langmuir isotherm was used to determine maximum
adsorption capacity of the HAP absorber and was found to be 123.45 mg/g. All calculated
thermodynamic parameters (ΔGº, ΔHº and ΔSº) clearly indicate sorption was thermodynamically
favourable, endothermic and spontaneous in nature.
Acknowledgements
Ms. Sridevi Brundavanam would like to acknowledge Murdoch University for providing a PhD Scholarship to
undertake the cadmium adsorption studies as part of her PhD project.
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American Journal of Materials Science 2015, 5(2): 31-40 DOI: 10.5923/j.materials.20150502.02
Kinetic and Adsorption Behaviour of Aqueous Fe2+, Cu2+ and Zn2+ Using a 30 nm Hydroxyapatite Based Powder
Synthesized via a Combined Ultrasound and Microwave Based Technique
Sridevi Brundavanam, Gérrard Eddy Jai Poinern*, Derek Fawcett
Murdoch Applied Nanotechnology Research Group, Department of Physics, Energy Studies and Nanotechnology, School of Engineering and Energy, Murdoch University, Murdoch, Australia
Abstract The present study reports the kinetic and absorption performance of novel nanometre scale hydroxyapatite (HAP) absorber synthesised from a combined ultrasound and microwave based technique for the removal of metal ions (Fe2+, Cu2+ and Zn2+) from aqueous solutions. After powder characterisation was carried out using XRD, SEM, EDS and FT-IR, batch adsorption studies were carried out. Kinetic studies established that Fe2+and Cu2+ ion adsorption tended to follow a pseudo-second order model, while Zn2+ ion adsorption tended to follow an intra-particle diffusion pattern. All three metal ion adsorption studies indicated an ion-exchange mechanism (metal ion → Ca2+) was a primary participant in the sorption process and was influenced by intra-particle diffusion. The Isotherm studies indicated the Langmuir isotherm modelled Fe2+ and Cu2+ ion adsorption, while the Freundlich isotherm was the better model for Zn2+ ion adsorption data. Maximum adsorption capacity of HAP determined via Langmuir isotherm was found to be 61.35 mg/g for Cu2+ ions, 55.25 mg/g for Fe2+ ions and 48.54 mg/g for Zn2+ ions. The study established HAP as an effective absorbent material for the removal of Fe2+, Cu2+ and Zn2+ loaded aqueous solutions.
Keywords Hydroxyapatite, Adsorption, Metal ion, Ultrasound, Microwaves
1. Introduction The presence of heavy metal contaminants in the
environment is a well-recognised problem that threatens the world today. Even in small concentrations, their high solubility in aquatic environments makes them extremely hazardous to living organisms [1, 2]. Once heavy metals enter the food chain their toxicity together with their tendency to accumulate beyond acceptable concentrations in the human body can result in serious health disorders [3]. Metals are generally considered heavy metals when their density exceeds 5 gcm-3 [4]. Metals and alloys with densities greater than 5 g cm-3 are widely used as biomaterials due to their superior mechanical properties. The majority of these metallic biomaterials are used in load-bearing implants such as bone plates, fixation pins and screws, hip and knee replacements. Metallic biomaterials are grouped into three distinct categories: 1) stainless steels; 2) cobalt-based alloys, and 3) titanium alloys. Individually,
* Corresponding author: [email protected] (Gérrard Eddy Jai Poinern) Published online at http://journal.sapub.org/materials Copyright © 2015 Scientific & Academic Publishing. All Rights Reserved
each category consists of a primary metallic biomaterial that is composed of a number of secondary metallic constituents. For example, the main constituents of stainless steels are iron (Fe), chromium (Cr) and nickel (Ni). Most of the metal constituents used in the three metallic biomaterial groups like Cr, cobalt (Co), Fe, molybdenum (Mo), Ni, and tungsten (W) are used as alloying elements to improve the properties of the main constituent element [5]. For example, Fe based stainless steels have Cr (17-20%) added to improve corrosion resistance and Mo (2-4%) is added to improve resistance to pitting corrosion [6]. However, metallic constituents such as Cr and Mo are toxic and can only be tolerated in very small amounts in the body. Studies have shown that leaching of metallic surgical implants can result in the release of toxic metallic ions into surrounding cells, blood vessels and tissues [7-9]. The release of metallic ions can only be tolerated in very minute amounts, as in larger amounts they induce an unfavourable inflammatory response that significantly reduces the biocompatibility of the implant [8]. Importantly, exposure to heavy metals can lead to serious health problems such as cancer, nervous system and organ damage. For example, both Cr and Ni are known carcinogens and long-term exposure can lead to the development and growth of cancer.
32 Sridevi Brundavanam et al.: Kinetic and Adsorption Behaviour of Aqueous Fe2+, Cu2+ and Zn2+ Using a 30 nm Hydroxyapatite Based Powder Synthesized via a Combined Ultrasound and Microwave Based Technique
Metallic biomaterials used in orthopaedic applications will generally experience load-induced stresses in the inner core of the implant. Whereas, the surface of the implant is exposed to the surrounding physiological environment in the body. The interaction between exposed surface and physiological environment is very important in soliciting a favourable biological response. Because of the importance that biocompatibility plays in delivering a successfully implantation procedure, there have numerous studies into improving interfacial properties between the implant surface and the physiological environment [9, 10]. To this end, bioactive coatings based on calcium phosphate ceramics have been extensively investigated and used to coat a variety of biomedical metallic implants. These coatings have been found to improve biocompatibility, promote implant attachment and restrain the release of metallic ions into the physiological environment [10, 12]. In particular, hydroxyapatite (HAP) is a thermodynamically stable mineral phase at physiological pH and has a hexagonal crystal structure composed of calcium phosphate groups. Both crystallographic and chemical studies reveal synthetic HAP is similar in chemical composition to the naturally occurring mineral phase found in human bone and teeth [13]. Synthetic HAP has the advantages of being biocompatible, non-toxic, and has enhanced bioactive properties towards bone cells and other body tissues [14].
Another interesting property of HAP results from its complex hexagonal structure. The structure provides effective high capacity absorbance for a variety of pharmaceutical products such as antibiotics, drugs, enzymes, hormones and steroids. The use of HAP as a slow and sustained release drug delivery platform has proven to be effective in the treatment of diseases such as osteomyelitis, osteoporosis and osseous cancer [15-18]. Likewise, HAP used in orthopaedic applications also has the potential to act as an adsorption matrix. In this case, adsorption results from a mass transfer process by which metallic particles are transferred from the physiological environment (liquid phase) to the HAP matrix, and become bound by physical and/or chemical interactions. HAP makes an attractive absorbent because of its advantageous surface properties such as its hydrophilic nature, surface charge, pH, porous structure and 2.6 P-OH surface groups per nm2 which act as sorption sites [19, 20]. Generally speaking there are three steps involved in sorption of metallic particles onto the HAP absorber: (1) transfer of metallic ion particles from the physiological environment to the absorber surface; (2) adsorption on the metallic ion particles onto the surface, and (3) transport of the metallic ion particles within the absorber matrix.
In the present work the adsorption of copper (Cu), iron (Fe), and zinc (Zn) ions from aqueous solutions onto a solid nanometre scale HAP powder in agitated batch absorber vessels was studied. The nanometre scale HAP powder was synthesized via a combined ultrasound and microwave based technique to produce a spherical particle with a mean diameter of 30 nm. The main goal of this study was to
examine the ability of the nanometre scale HAP powder to remove these ions from an aqueous solution and therefore evaluate its potential to store heavy metallic ions. The source of these metallic ions could arise from the corrosion and wear of surgical implants or by metabolic uptake via nutritional intake, drinking water and inhalation. The synthesized HAP powders were characterized using X-ray diffraction (XRD) spectroscopy, Scanning electron microscopy (SEM), Energy Dispersive Spectroscopy (EDS) and Fourier Transform Infrared Spectroscopy (FTIR). The metal ion adsorption capacity of the powders were investigated via the removal of Cu, Fe, and Zn ions from aqueous solutions using a batch equilibrium procedure. The kinetic behaviour of metal ion adsorption was investigated using Lagergren’s pseudo-first order, McKay & Ho’s pseudo-second order and intra-particle diffusion models. Furthermore, Langmuir and Freundlich adsorption isotherms were used to model the experimental data.
2. Materials and Methods 2.1. Materials
Adsorption experiments were conducted using three different metal salts. The first, FeCl2.4H2O was supplied by Chem-Supply (Australia). The second, CuCl2.2H2O was supplied by Sigma-Aldrich (United States) and the third, ZnCl2 was supplied by Scharlau (Barcelona, Spain). Each respective metal salt was dissolved in an aqueous solution to make up a 1000 ppm stock solution and lower solution concentrations were produced by successive dilution of the stock solution. HAP powders were synthesized from high purity calcium nitrate tetra-hydrate [Ca(NO3)2.4H2O] and potassium di-hydrogen phosphate [KH2PO4], while solution pH was controlled by the addition of ammonium hydroxide [NH4OH]. All chemicals used to synthesize HAP were supplied by Chem-Supply (Australia). During HAP synthesis an Ultrasound Processor [UP50H: 50 W, 30 kHz, MS7 Sonotrode (7mm diameter, 80 mm length)] supplied by Hielscher Ultrasound Technology was used to deliver the ultrasound irradiation. All aqueous solutions were made using Milli-Q® water (18.3 MΩ cm-1) produced by an ultrapure water system (Barnstead Ultrapure Water System D11931; Thermo Scientific, Dubuque, IA).
2.2. Synthesis of Nanometres Scale Hydroxyapatite
Ultrasonic and microwave processing are two techniques that significantly influence and enhance the properties of materials. Ultrasound produces extremely rapid pressure and temperature variations that promote both physical effects and chemical reactions that directly influence particle size and morphology during synthesis. Microwave processing has the advantages of providing controlled volumetric heating, lower energy consumption, reduced reaction times, increased product yields, and improved material properties when compared to conventional heating methods [21]. A detailed synthesis procedure developed by
American Journal of Materials Science 2015, 5(2): 31-40 33
the authors is given elsewhere [21, 22]. But a brief description is give here in the interest of completeness and begins by adding a 40 mL solution of 0.32 M calcium nitrate tetra-hydrate into a small glass beaker. The pH of the solution is then adjusted to 9.0 by adding approximately 2.5 mL of ammonium hydroxide. The solution was then subjected to 50 W of ultrasound irradiation for 1 h set and maximum amplitude. The second hour of ultrasound treatment included slowly adding 60 mL of 0.19 M potassium di-hydrogen phosphate solution. During the addition of potassium di-hydrogen phosphate the solution pH was maintained at 9.0 and the Calcium/Phosphate [Ca/P] ratio was maintained at 1.67. After ultrasonic treatment, the solution was centrifuged (15,000 g) for 20 minutes at room temperature to produce a solid white precipitate. The precipitate was collected, washed and centrifuged for a further 10 minutes. At the end of the second centrifugation, the precipitate was placed into a fused silica crucible and then loaded into a domestic microwave oven for thermal treatment [Set at 100% power for 40 minutes: 1100W at 2450 MHz-LG® Australia] [21]. The resultant agglomerated powder was collected and then subjected to ball milling until an ultrafine nanometre scale HAP powder was produced.
2.3. Advanced Characterisation Techniques
2.3.1. X-ray Diffraction (XRD) Spectroscopy
XRD spectroscopy was used to study the synthesized powders, with the XRD patterns being used to identify the crystalline size and phases present. Spectroscopy data was recorded at room temperature, using a Siemens D500 series diffractometer [Cu Kα = 1.5406 Å radiation source] operating at 40 kV and 30 mA. The diffraction patterns were collected over a 2θ range of 20° to 60° with an incremental step size of 0.04° using flat plane geometry with 2 second acquisition time for each scan. The scan data was then used in conjunction with the Debye-Scherrer equation [Equation 1] to determine the crystalline size of each sample.
2.3.2. Scanning Electron Microscopy (SEM) and Energy Dispersive Spectroscopy (EDS)
The SEM technique was used to examine the size, shape and morphological features of absorbents before and after the batch adsorption studies. All micrographs were taken using a JCM-6000, NeoScopeTM with attached energy dispersive X-ray spectroscopy. Samples were mounted on individual substrate holders using carbon adhesive tape before being sputter coated with a 2 nm layer of gold to prevent charge build up using a Cressington 208HR High Resolution Sputter coater.
2.3.3. Fourier Transform Infrared Spectroscopy (FT-IR)
FT-IR spectroscopy was used to identify species and functional groups present in the HAP absorber before and after metal adsorption studies. Samples were examined
using a Perkin–Elmer Frontier FT-IR spectrometer with Universal Single bounce Diamond ATR attachment. Both FT-IR spectra were recorded in the range from 525 to 4000 cm−1 in steps of 4 cm-1.
2.4. Batch Adsorption Studies
All adsorption experiments were carried out using the batch equilibrium technique. During the experimental work the adsorption capacity of the HAP absorber for three metal ions, namely, Cu2+, Fe2+ and Zn2+ were investigated. The study also examined the influence of the initial metal ion concentration and contact time at 298ºC. The influence of contact time on metal ion adsorption on the absorber was examined using aqueous solutions containing 300 mg/L of metal ions (300 ppm) prepared from the three metal chlorides. HAP samples (0.1 g) were taken from the stock solution (1g/L) were added to each of the 300 ppm metal chloride solutions. The magnetic stirring speed for each metal suspension were set to 400 rpm, while the temperature of each suspension was maintained at 298 ± 1 K. In addition, the initial and final pH of the suspensions were recorded. Sample volumes were taken from the suspension during the mixing process at pre-determined time intervals (10, 20, 30, 40, 60, 90, 120, 180, 240 and 300 min) so that metal ion concentration in the respective solution could be measured. Once a sample volume was taken, an equivalent volume was added to maintain testing solution volume. Each solution volume was filtered using a Whatman® 0.22μm membrane syringe filter before being centrifuged at 15,000 g for 20 minutes. The concentration of Cu2+, Fe2+, Zn2+ and Ca2+ ions in the sample solutions were determined using atomic absorption spectroscopy (AAS). The instrument used was a Varian SpectrAA50 (Victoria, Australia) flame atomic absorption spectrometer operated in accordance with the manufacturer’s recommendations. Elemental analysis of Cu, Fe and Zn was carried out using an air–acetylene flame and Ca analysis was carried using a nitrous oxide (N2O)–acetylene flame. Sample aspiration flow rate used was 5 mL min−1. During analysis hollow cathode lamps of Fe, Cu, Zn and Ca (Varian) were used. While Fe, Cu, Zn and Ca concentration were measured at the wavelengths of 248.3, 324.8, 213.9 and 422.7 nm respectively. In addition, the influence of initial metal ion concentration was studied by first preparing a series of Schott reagent bottles containing 100 mL aqueous solutions consisting of varying concentrations of metal ions (100, 150, 200, 250 and 300 mg/L). All initial metal ion concentration experiments were carried out in triplicate. The data collected from the adsorption experiments were then used in the subsequent kinetic and adsorption isotherm modelling studies.
3. Results and Discussions 3.1. XRD Spectroscopy, SEM and EDS Analysis
XRD spectroscopy was carried out on all HAP powders
34 Sridevi Brundavanam et al.: Kinetic and Adsorption Behaviour of Aqueous Fe2+, Cu2+ and Zn2+ Using a 30 nm Hydroxyapatite Based Powder Synthesized via a Combined Ultrasound and Microwave Based Technique
before and after adsorption studies. The resultant XRD patterns were used to determine crystalline size and phases present in the samples. An XRD pattern of a typical pure powder sample before adsorption testing is presented in Figure 1 and is indicated by the purple pattern. Examination of the pure HAP sample pattern reveals the presence of peaks that coincide with the known phases of pure HAP and is consistent with the phases listed in the ICDD database. The main (h k l) indices found in pure HAP, namely (002), (211), (112), (300), (202), (310), (222), (213) and (004) can be seen in the pattern for the synthesized HAP powder. In addition, the pattern showed no evidence of non-HAP phases. The crystalline size, t (hkl), of each sample was calculated from the respective XRD patterns using the Debye-Scherrer equation [23-25]
(1)
where, λ is the wavelength of the monochromatic X-ray beam, B is the Full Width at Half Maximum (FWHM) of the peak at the maximum intensity, θ(hkl) is the peak diffraction angle that satisfies Bragg’s law for the (h k l) plane and t(hkl) is the crystallite size. The mean crystallite size calculated from the (002) reflection peak for pure HAP sample was found to be 30 nm.
After adsorption studies, XRD spectroscopy was carried out on the metal ion loaded samples and representative XRD patterns are presented in Figure 1. XRD analysis
reveals that there was no major changes in peak locations and powder size resulting from metal ion adsorption, as seen in Figure 1. To investigate this result further, SEM and EDS analysis was carried out on all samples. SEM analysis revealed that all powder samples displayed the same agglomerated spherical and granular morphology.
Figure 1. XRD patterns of powder samples before and after adsorption studies. The lower (Purple) pattern is a typical unloaded HAP powder sample, while the second (Red) a Fe2+ loaded sample. This is followed by a Cu2+ (Green) loaded sample and finally a Zn2+ (Blue) loaded sample
Figure 2. (a) SEM of a pure HAP powder; (b) EDS of pure HAP sample with a Ca:P ratio of 1.66; (c) SEM micrograph of an Fe2+ loaded sample, and (d) EDS spectrum confirming the presence of Fe2+ ions in the powder sample
)()( θcos
9.0
hklhkl B
t λ=
American Journal of Materials Science 2015, 5(2): 31-40 35
Figure 3. (a) SEM micrograph of Cu2+ loaded sample; (b) EDS spectrum of Cu2+ loaded sample; (c) SEM micrograph of an Zn2+ loaded sample, and (d) EDS spectrum confirming the presence of Zn2+ ions
Figure 2 (b) presents a typical EDS spectrum of an unloaded sample showing dominant peaks corresponding to Ca, P, and O, and confirms the chemical composition of HAP. Analysis of the EDS data revealed a Ca:P ratio of 1.66, which was very close to the ideal value of 1.67 normally associated with HAP. Also present was a minor peak that resulted from the carbon adhesive tape used to attach the sample to the SEM stub. Figure 2 (c) presents an SEM micrograph of a representative Fe2+ loaded sample that also displays the highly agglomerated nature of the synthesized powder. The accompanying EDS spectrum confirms the presence of Fe2+ ions present in the sample. The EDS also shows that there has been a reduction in the intensity of the Ca peak and suggests Fe2+ ions have replaced Ca2+ ions in the lattice. Studies have shown that divalent metal cations have a strong selectivity towards the Ca2+ in the HAP lattice via an ion-exchange mechanism [26, 27]. A similar trend can be seen for Cu2+ and Zn2+ ion load samples as seen in Figure 3. However, the reduction in the Ca peak for the Cu2+ and Zn2+ ion load samples is not as pronounced as the Fe2+ loaded sample seen in Figure 2 (d). The ion-exchange mechanism between the respective metal ion and Ca2+ can be represented by the equivalent molar exchange in the general HAP formula Ca10−x Mx(PO4)6 (OH)2, where M represents the respective metal ion and x can vary from 0 to 10 depending on experimental parameters and reaction time. Studies reported in the literature have detected very small XRD shifts associated
with the ion-exchange mechanism for divalent metal ions [26, 28]. Inspection of the XRD patterns of representative metal ion loaded samples presented in Figure 1 do not reveal any of these very small XRD shifts.
3.2. FT-IR Spectroscopy Studies
FT-IR spectroscopy was used to detect species and functional groups associated with peaks seen in sample spectra taken before and after adsorption studies. Figure 4 presents the results of a FT-IR spectroscopy study of HAP powder samples taken before and after metal ion adsorption studies. Starting from the right hand side of Figure 4 with a typical synthesized HAP powder sample (purple) prior to adsorption studies. The first three peaks encountered are 561 cm-1, 600 and 631cm-1 that are associated with v4 vibrations of the O-P-O modes. The peak located at 827 cm-1 indicates the presence of carbonates in the sample and is a consequence of atmospheric carbon dioxide interacting with HAP precursors during synthesis and has been reported in the literature by other researchers [29, 30]. The small peak located at 963 cm-1 is associated with v1 symmetric stretching vibrations associated with a P-O mode. While the much larger peak at 1026 cm-1 and the smaller peak located at 1090 cm-1 correspond to PO4
3- functional groups. While the peaks located at 1373 cm-1 and 1643 cm-1 correspond to CO3
2- functional groups. While the weak peak located at 3470 cm-1 corresponds to OH- ion vibrations in the HAP crystal lattice. The three remaining spectra are
36 Sridevi Brundavanam et al.: Kinetic and Adsorption Behaviour of Aqueous Fe2+, Cu2+ and Zn2+ Using a 30 nm Hydroxyapatite Based Powder Synthesized via a Combined Ultrasound and Microwave Based Technique
representative Fe2+, Cu2+ and Zn2+ loaded powder samples taken after adsorption testing. The results for the metal ion loaded samples are very much the same as the pre-adsorption sample, except that both 872 cm-1 and 1373 cm-1 peak intensities have vanished from all spectra. Both the missing peaks correspond to CO3
2- functional groups and suggest that the low pH of the batch adsorption technique resulted in the loss of the carbonates. For example, equilibrium pH of the Fe2+ ion and HAP solution pH was 3.2, while Cu2+ was 3.8 and Zn2+ was 4.4.
Figure 4. FT-IR Spectroscopy analysis of synthesized pure nanometre scale hydroxyapatite powder and metal ion loaded samples after adsorption studies
3.3. Adsorption Kinetics
A proper understanding of adsorption kinetics is important since sufficient residence time on the absorber surface is needed to complete the adsorption reaction. This is reflected in the metal ion concentration profiles encountered during the batch adsorption studies and presented in Figure 5. Inspection of the profiles reveals there is an initial rapid uptake of metal ions, but as time progresses the uptake slows and around the 150 minutes mark no further adsorption takes place. The quantity of metal ions adsorbed at equilibrium time (qe) was determined by equation 2 below:
𝑞𝑞𝑒𝑒 = (𝐶𝐶𝑜𝑜 − 𝐶𝐶𝑒𝑒) 𝑉𝑉𝑚𝑚
(2)
where Co and Ce are the initial and equilibrium concentrations (mg/L) of metal ions in solution, V is solution volume (L) and m is absorber mass (g) used during the experiments. Three kinetic models were used to examine the experimental data. The first was the Lagergren pseudo-first order law [31] that is defined by equation 3 below:
log (𝑞𝑞𝑒𝑒 − 𝑞𝑞𝑡𝑡) = log 𝑞𝑞𝑒𝑒 − 𝑘𝑘12.303
𝑡𝑡 (3)
where, qt (mg/g) is adsorption at time t and k1(min-1) is the pseudo-first order adsorption rate constant. The second model used was the McKay and Ho’s pseudo-second-order law [32] and is presented in equation 4:
𝑡𝑡𝑞𝑞𝑡𝑡
= 1𝑘𝑘2𝑞𝑞𝑒𝑒2
+ 1𝑞𝑞𝑒𝑒
𝑡𝑡 (4)
where k2 (g/min.mg) is the pseudo-second-order rate constant for adsorption. The third model used was a multi-step intra-particle diffusion model proposed by Weber and Morris [33] and is defined by equation (5) below:
𝑞𝑞𝑡𝑡 = 𝑘𝑘𝑝𝑝 𝑡𝑡12 + 𝐶𝐶 (5)
where C is the intercept that provides the ideal boundary layer thickness and kp is the intra-particle diffusion rate constant (mg/g.min1/2 g).
Figure 5. Metal ion concentration profiles during batch adsorption studies over a period of 300 minutes using a nanometre scale HAP powder as absorber
Figure 6 (a) presents kinetic data plotted using the Lagergren pseudo-first order equation (3), Figure 6 (b) displays the kinetic data plotted using the McKay and Ho’s pseudo-second-order equation (4) and Figure 6 (c) represents the multi-gradient line data of intra-particle diffusion model (5). Inspection of Figures 6 (a) and (b) reveals that both the pseudo-first order and pseudo-second-order models gave reasonable representation of the data obtained. In spite of the good initial agreement, further analysis reveals the pseudo-second order model gives a better representation of the data for Fe2+ and Cu2+ ions with a slightly higher correlation coefficient (R2) as seen in Table 1. However, in the case of Zn2+ ions, the intra-particle diffusion model gave a better representation of the data as seen in Figure 6 (c) and Table 1.
3.4. Adsorption Isotherms
There are two widely used equilibrium equations for modelling equilibrium data obtained from adsorption systems. The first model is the Freundlich and is purely an empirical equation that takes into account surface heterogeneity, the exponential distribution of active adsorption sites and their respective energies over a wide range of concentrations. The second model is the Langmuir equation and unlike the Freundlich, it assumes maximum adsorption occurs when the entire surface of the absorber is
American Journal of Materials Science 2015, 5(2): 31-40 37
covered by a monolayer of adsorbate. The equilibrium data for metal ions in solution for initial concentrations consisting of 100, 150, 200, 250 and 300 mg/L at constant temperature of 298 K, 1 g/L absorbent dose and a contact time of 300 minutes were analysed using Freundlich and Langmuir isotherms. The linear form of the Freundlich isotherm used for modelling the data is expressed by in equation (6):
log 𝑞𝑞𝑒𝑒 = log𝑘𝑘𝐹𝐹 + 1𝑛𝑛
log𝐶𝐶𝑒𝑒 (6)
Table 1. A comparison between the pseudo kinetic (first and second order) rate constants and intra-particle kinetic diffusion constants
Pseudo-first-order kinetic model
Metal ion Temperature (K) k1 (min-1) qe
(mg/g) R2
Fe2+ 298 33.92 55.24 0.9904
Cu2+ 298 19.99 62.50 0.9805
Zn2+ 298 115.53 57.14 0.9875
Pseudo-second-order kinetic model
Metal ion Temperature (K)
k2
(g/mg.min) qe
(mg/g) R2
Fe2+ 298 6.84 x 10-4 52.91 0.9951
Cu2+ 298 12.15 x 10-4 59.17 0.9980
Zn2+ 298 3.85 x10-4 42.01 0.9734
Intra-particle kinetic diffusion model
Metal ion Temperature (K)
kp (g/mg.min)
C (mg/g) R2
Fe2+ 298 4.29 1.12 0.9934
Cu2+ 298 4.17 13.17 0.9080
Zn2+ 298 3.77 -7.98 0.9995
where kF and n are Freundlich parameters related to the extent of adsorption and the intensity of adsorption respectively. Both kF and n parameters were determined via plotting log qe versus log Ce , the results of which are presented in Figure 7 (a). The linear form of the Langmuir isotherm used for modelling the data is expressed by equation (7):
𝐶𝐶𝑒𝑒𝑞𝑞𝑒𝑒
= 1𝑄𝑄𝑚𝑚 𝑏𝑏
+ 𝐶𝐶𝑒𝑒𝑄𝑄𝑚𝑚
(7)
where, Qm (mg/g) is the monolayer adsorption capacity, b (L/g) is the Langmuir constant that is related to the free energy of adsorption, while Ce (mg/L) is the equilibrium concentration of adsorbate in solution and qe (mg/g) is the concentration of adsorbate on the surface of absorber.
Both Qmax and b parameters were determined via plotting Ce/qe versus Ce over the entire metal ion concentration range as seen in Figure 7 (b). The parameters reflect the surface properties and affinity of the metal ions for the nanometre scale HAP absorber. Both Langmuir and Freundlich plots exhibited good linear fits and were able to provide slope parameters and intercepts. The results of the respective isotherm analysis are listed in Table 2. For the Freundlich isotherm, if n falls within the range of 1 to 10 it indicates the
adsorption process is favourable. In this study, n for all metal ions fell within this range and confirmed that adsorption was favourable. Furthermore, based on the coefficient of correlation (R2) values, the Langmuir isotherm provided the best fit for Fe2+ and Cu2+ experimental data. While the Freundlich isothermal provided the best fit for Zn2+
experimental data. The maximum monolayer adsorption capacity (Qm) for respective metal ions was calculated using the Langmuir equation and revealed the maximum adsorption capacity occurred for Cu2+ ions (61.35 mg/g). This was followed by Fe2+ ions (55.25 mg/g) and then Zn2+ ions (48.54 mg/g). The study also found that adsorption capacity was the result of increasing equilibrium metal ion concentrations in solution. The increased concentrations were able to increase the numbers of metal ions at the absorber surface and enhance the probability of adsorption.
Figure 6. Metal ion adsorption data modelled using three kinetic models: (a) Lagergren’s pseudo-first order law; (b) McKay and Ho’s pseudo-second-order law, and (c) the intra-particle diffusion model
38 Sridevi Brundavanam et al.: Kinetic and Adsorption Behaviour of Aqueous Fe2+, Cu2+ and Zn2+ Using a 30 nm Hydroxyapatite Based Powder Synthesized via a Combined Ultrasound and Microwave Based Technique
Table 2. Comparisons between Freundlich and Langmuir adsorption isotherm constants for metal ions (Fe2+, Cu2+, Zn2+) onto nanometre scale hydroxyapatite at 298 K
Freundlich adsorption isotherm constants
Metal ion
Temperature (K) kf (mg/g) n R2
Fe2+ 298 19.02 5.685 0.9986
Cu2+ 298 21.17 5.640 0.9975
Zn2+ 298 3.53 2.448 0.9919
Langmuir adsorption isotherm constants
Metal ion
Temperature (K)
Qmax (mg/g) b (L/mg) R2
Fe2+ 298 55.25 0.036 0.9992
Cu2+ 298 61.35 0.038 0.9993
Zn2+ 298 48.54 0.009 0.9817
Figure 7. Linear fits of experimental data using (a) Freundlich and (b) Langmuir isotherms
Table 3 presents a comparison between the results of this study and with those reported in the literature for selected adsorbents. Comparing the adsorption capacity of HAP powder used in this study with the different adsorbents listed in Table 3 reveals the powder is an effective adsorbent for the removal of Fe2+, Cu2+ and Zn2+ ions from aqueous solutions. Because of the relatively low-cost of the starting materials and the optimised ultrasound/microwave route used, the resultant HAP powder has the potential to be
a cost-effective adsorbent for the removal of Fe2+, Cu2+ and Zn2+ ions from aqueous solutions.
Table 3. Adsorption capacities for sorption of Fe2+, Cu2+ and Zn2+ ions from aqueous solutions by various adsorbents
Adsorbent Q (mg/g) Reference
Fe2+ Cu2+ Zn2+
Activated carbon derived from Cicer arietinum - 17.77 19.93 34
Bentonite - 14.10 21.1 35
Bone Charcoal - 45.80 35.15 36
Cow Bone Charcoal 31.43 35.44 - 37
Egg Shell - 34.48 35.71 38
Hydroxyapatite - 36.9 - 39
Hydroxyapatite (30 nm) 55.25 61.35 48.54 This Study
4. Conclusions The results of the present study have revealed that a
nanometre scale HAP powder synthesized using a combined ultrasound and microwave based technique was capable of producing a highly crystalline powder consisting of spherical particle morphology with a mean particle size of 30 nm. The powder was found to be an effective adsorbent for the removal of metal ions such as Fe2+, Cu2+ and Zn2+ from aqueous solutions. Kinetic studies revealed the sorption process for Fe2+and Cu2+ ions closely followed pseudo-second order kinetics. While the slightly higher correlation coefficient (R2) of 0.9995 for Zn2+ ions indicated that the Intra-particle kinetic diffusion kinetics was more suited to modelling the experimental data. Overall, the sorption performance was found to be a function of initial metal ion concentration. With initial uptake rate of metal ions being high compared to the much lower uptake rates occurring in the later part of the absorption period. The ion-exchange mechanism (metal ion → Ca2+) was seen to be a participant in the sorption process which was influenced by intra-particle diffusion. Isotherm studies indicated the Langmuir isotherm modelled Fe2+ and Cu2+ ion adsorption data better than the Freundlich isotherm. However, the Freundlich isotherm modelled the Zn2+ ion adsorption data. The Langmuir isotherm was used to determine maximum adsorption capacity of the HAP absorber for each metal ion. The study found the metal ion with the largest maximum adsorption capacity was Cu2+ ions (61.35 mg/g). This was followed by Fe2+ ions (55.25 mg/g) and then Zn2+ ions (48.54 mg/g).
ACKNOWLEDGEMENTS Mrs Sridevi Brundavanam would like to acknowledge
Murdoch University for providing a Postgraduate Scholarship to undertake this hydroxyapatite synthesis study as part of her PhD project.
American Journal of Materials Science 2015, 5(2): 31-40 39
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52
4.3. Chapter Summary
This chapter consisted of two case studies that examined the chemical adsorption
capabilities of HAP and its ability to function like the mineral phase of bone tissue. The
prepared nanometre scale HAP powders were found to be an effective adsorbent for the
removal of metal ions (Cd2+
, Fe2+
, Cu2+
and Zn2+
) from aqueous solutions. The sorption
performance of the metal ions was found to be a function of initial metal concentration,
temperature and solution pH. Furthermore, metal ion removal was found to improve
with increases in these parameters for specific contact times. Kinetic studies were also
carried out and revealed that during sorption, the initial uptake rate of metal ions was
high, but this was followed by a much lower uptake rate. The ion-exchange mechanism
(Metal ion → M2+
) was clearly identified as a participant in the sorption process.
53
Chapter 5 - Growth of flower-like brushite structures on magnesium
substrates and their subsequent low temperature transformation to
hydroxyapatite.
5.1. Overview and author contributions
Chapter five addresses two components outlined in aims 3 and 4 of this research project.
The first component, listed in aim 3 was to develop and use a direct chemical
immersion technique to form calcium phosphate coating on magnesium substrates. The
second component was to conduct corrosion studies to investigate the degradation
behaviour of various coating types in phosphate buffer saline (PBS) solution and
Ringer’s solution at 37 ºC and at a pH of 7.4 to simulate the body’s physiological fluid
environment. Case study 4 is composed of two parts. The first part developed a
straightforward and cost effective chemical immersion technique that deposited di-
calcium phosphate dihydrate or brushite [DCPD; CaHPO4.2H2O] coatings on Mg
substrates. In the second half of the study the DCPD coatings were converted into
hydroxyapatite [(HAP); Ca10(OH)2 (PO4)6] via immersion in a sodium hydroxide
solution bath at 80 ºC for 2h. The size, morphology and crystallinity of both DCPD and
HAP coating were examined using X-ray diffraction (XRD) spectroscopy, Field
emission scanning electron microscopy (FESEM), Energy Dispersive Spectroscopy
(EDS) and Fourier Transform Infrared spectroscopy (FT-IR). The second part of the
research was to investigate the degradation behaviour of pure Mg substrates and the
various substrate coatings. The degradation behaviour was evaluated in two electrolytes,
the first was phosphate buffer saline (PBS) solution and the second was Ringer’s
solution. All degradation studies were carried out at human body temperature (37 ºC).
54
S. Brundavanam conducted all chemical immersion studies, low temperature
transformation procedures and performed advanced characterisation techniques [(XRD),
(SEM), (EDS) and (FT-IR)] to fully investigate the physical and chemical properties of
pure Mg substrates and substrate coatings. S. Brundavanam also carried out all
electrochemical measurements associated with substrate degradation studies with the
assistance of Associate Professor Gamini Senanayake. S. Brundavanam actively
participated in the subsequent analysis of the data from chemical immersion studies and
advanced characterisation studies. S. Brundavanam carried out all studies forming Case
Study 4 under the supervision of G.E.J. Poinern. During the studies D. Fawcett
(MANRG - Research Fellow) assisted S. Brundavanam in addressing the various
technical aspects of the research. S. Brundavanam also contributed to writing the
research article and assisted with editorial changes recommended by reviewers for the
peer-reviewed research article that formed Case Studies 4. All authors provided feed-
back during the preparation of the paper.
5.2. Published Research Article:
Case Study 4
Sridevi Brundavanam, Gérrard Eddy Jai Poinern, Derek Fawcett. Growth of
flower-like brushite structures on magnesium substrates and their subsequent low
temperature transformation to hydroxyapatite. American Journal of Biomedical
Engineering. 2014; 4(4): 79-87.
American Journal of Biomedical Engineering 2014, 4(4): 79-87 DOI: 10.5923/j.ajbe.20140404.02
Growth of Flower-Like Brushite Structures on Magnesium Substrates and Their Subsequent Low Temperature Transformation to Hydroxyapatite
Sridevi Brundavanam, Gérrard Eddy Jai Poinern*, Derek Fawcett
Department of Physics, Energy Studies and Nanotechnology, School of Engineering and Energy, Murdoch University, Murdoch, Australia
Abstract Dicalcium phosphate dihydrate (DCPD) Brushite coatings composed of flower-like structures were formed on magnesium substrates via a straightforward chemical immersion technique in order to slow down the corrosion rate of the metallic substrates. Moreover, the synthesised DCPD coatings were also converted to hydroxyapatite (HAP) coating using a low-temperature hydrothermal process to further investigate their ability to reduce the corrosion rate of the substrates in phosphate buffer saline (PBS) and Ringer’s solutions. Degradation studies found DCPD coatings were capable of providing the most significant reduction in the corrosion rate of around 0.100 mm/yr compared to 3.828 mm/yr for the uncoated substrates soaked in Ringer’s solution at 37ºC.
Keywords Magnesium, DCPD, Hydroxyapatite, Coating, Simulated Biodegradation
1. Introduction Ideally, orthopaedic and dental devices must have a
surface chemistry and physical structure that is biological compatible. In addition, they must also induce a positive osteogenic response and avoid any unfavourable immune responses [1]. Biodegradable implant materials currently in use and under development offer many attractive features for a number of clinical applications such as bone fixation, controlled release of pharmaceuticals, endovascular stents and orthopaedic implants [2-4]. The use of biodegradable materials in orthopaedic surgical procedures is of particular interest, since the function of many implants usually comes to an end when tissue regeneration and healing has taken place [3, 5]. The majority of these implants are biologically inert and engineered from metallic materials such as cobalt-chromium based alloys, stainless steels and titanium alloys. Metallic materials are very appealing for load bearing applications due to their ductility, high strength, fracture toughness and anticorrosion properties [6-8]. Typically, many of these implants are only needed for a short period of time to provide the necessary structural and mechanical support during tissue regeneration. However, to remove the implant after tissue regeneration requires a second surgical procedure. The additional surgical procedure is both costly and significantly increases the risk
* Corresponding author: [email protected] (Gérrard Eddy Jai Poinern) Published online at http://journal.sapub.org/ajbe Copyright © 2014 Scientific & Academic Publishing. All Rights Reserved
of infection and scarring of the patient [9]. Alternatively, the implant can be left in situ and as a result a number of potentially detrimental effects can take place such as corrosion of the implant material itself, inflammatory responses, stress-shielding and subsequent weakening of surrounding tissues. For example, there can be a significant release of toxic metallic ions such as chromium, cobalt and nickel during biological corrosion and mechanical wear. The production of these toxic metallic ions immediately solicits an unfavourable inflammatory response from the body’s immune system. The unfavourable immune response significantly reduces the biocompatibility of the implant and often leads to secondary revision surgery [10-12]. The other major disadvantage of metallic implants is their superior mechanical properties that are often many times greater than those of natural bone tissues. For instance, cobalt-chromium based alloys can have an elastic modulus that is ten times greater than that of bone and titanium based alloys are typically five times greater [6, 13, 14]. The significant difference between the mechanical properties of the implant and the surrounding bone creates a stress-shielding effect that causes bone resorption and subsequent implant failure. Furthermore, metallic implants are biologically inert and do not biologically or chemically interact with the surrounding tissues. The lack of interaction results in very little interfacial bonding or osteointegration taking place [15]. One currently used technique to improve the osteointegration of metallic implants involves coating them with a bioactive material such as calcium phosphate [16].
An attractive alternative to conventional metallic
80 Sridevi Brundavanam et al.: Growth of Flower-Like Brushite Structures on Magnesium Substrates and Their Subsequent Low Temperature Transformation to Hydroxyapatite
implants is to develop a biologically degradable implant that achieves complete dissolution by the end of the tissue regeneration period [17]. Some successful outcomes have been achieved using biodegradable polymers in applications such as sutures, bone and dental cements, bone grafting materials, plates, screws, pins and fixation devices [18-24]. However, their low mechanical strength means they are only suitable for low-load bearing applications and soft tissue reconstruction. On the other hand magnesium (Mg) is a metallic material with the potential to overcome the limitations of conventional metallic implant materials and degradable polymeric materials. Magnesium offers a number of attractive features such as biocompatibility, biodegradability and mechanical properties similar to bone [25]. The density of Mg is 1.74 g/cm3 at 20 and is slightly less than bone which ranges from 1.8 to 2.1 g/cm3. There is also a close similarity between the elastic modulus of Mg (45 GPa) and bone which varies from 40 to 57 GPa [26, 27]. The close similarity of its mechanical properties and its favourable biocompatibility makes Mg a promising material for the development of biodegradable orthopaedic implants [25, 28].
Despite Mg many advantages, but its poor corrosion resistance in chloride rich body fluids (pH ranges between 7.4 and 7.6) has severely limited its use in medical applications. The rapid corrosion rate results in two fundamental problems. The first involves the rapid formation of subcutaneous hydrogen gas bubbles that appear during the first week after surgery [29, 30]. However, the bubbles can be drawn off using a subcutaneous needle [31]. The second problem results in the loss of mechanical integrity between the implant and the surrounding bone tissue that prevents effective tissue regeneration. However, this vulnerability to corrosion can be considered as an advantage when designing a biodegradable implant. For example, controlled degradation of the implant will allow natural bone tissues to regenerate and replace the implant [32]. Generally, Mg corrodes in an aqueous environments according to the following equations. The anodic reaction can be explained by the partial reaction expressed in equation (1). While the partial reaction occurring at the cathode can be expressed by equation (2).
Anodic reaction: Mg → Mg2- + 2e- (1) Cathodic reaction: 2H2O + 2e- → 2OH- + H2 (2)
The complete corrosion process is presented in equation (3). However, corrosion occurring in the body environment is not straight forward and is complicated by the influence of factors such as: 1) the pH of body fluids; 2) variations in the pH value; 3) concentration of ions; 4) the presence of proteins and protein adsorption on the surface of the implant material; and 5) the influence of the surrounding tissues [25, 33, 34].
Mg (s) + 2H2O (l) → Mg (OH)2 (s) + H2 (g) (3) The most significant by-products produced during
degradation are hydrogen gas and Mg ions. Studies by Song
have suggested that hydrogen levels of around 0.01 ml/cm2/ day does not constitute a serious threat to body tissues [35]. While other studies have shown the release of Mg ions can promote cellular adhesion and cellular differentiation of bone cells [36, 37]. Another positive effect resulting from the release of Mg ions is their antibacterial properties that have the potential to prevent biological film formation on implant surfaces [38]. The results of the studies clearly indicate reducing the corrosion rate will reduce the formation levels of Mg ions and hydrogen to acceptable levels and make Mg an ideal biodegradable implant material. One effective method of reducing the corrosion rate is to coat it with a non-corrosive protective layer. Therefore, developing an effective biocompatible coatings capable of regulating the corrosion rate is essential for the development of a biodegradable Mg implants.
Calcium phosphates (CaP) have been widely used to coat metal implants because of their excellent biocompatibility, non-toxicity, bioactivity and bone inductivity properties. Several techniques have been used to deposit CaP coating on Mg and other metal substrate materials such as anodization [39], biomimetic methods [40], electro-less deposition [41], electro-deposition [42, 43], ion-beam-assisted deposition [44], chemical [45, 46] and hydrothermal [47]. Many of these techniques require complex equipment, multiple processing steps and high temperature treatments to produce an effective substrate coating.
In this study a straightforward and cost effective chemical immersion technique was used to deposit di-calcium phosphate dehydrate or brushite [DCPD; CaHPO4.2H2O] coatings on Mg substrates. The DCPD coatings were subsequently converted into hydroxyapatite [(HAP); Ca10(OH)2 (PO4)6] via immersion in a sodium hydroxide solution at 80 ºC for 2h. The size, morphology and crystallinity of both DCPD and HAP coating were examined using X-ray diffraction (XRD) spectroscopy, Field emission scanning electron microscopy (FESEM), Energy Dispersive Spectroscopy (EDS) and Fourier Transform Infrared spectroscopy (FT-IR). The degradation behaviour of Mg and Mg coated substrates were evaluated in phosphate buffer saline (PBS) solution and Ringer’s solution at body temperature (37ºC).
2. Materials and Methods 2.1. Materials
All chemicals used in this work were supplied by Chem-Supply (Australia) and all aqueous solutions were made using Milli-Q® water (18.3 MΩ cm-1) produced by an ultrapure water system (Barnstead Ultrapure Water System D11931; Thermo Scientific, Dubuque, IA).
2.2. Mg Substrate Pre-treatment
Mg (99.9% pure) sheets were cut into rectangular strips 40 mm in length, 3 mm in width and 0.15 mm in thickness.
American Journal of Biomedical Engineering 2014, 4(4): 79-87 81
The strips were polished using 1200 grit silicon carbide (SiC) paper to remove all surface oxides and contaminants. After polishing, the strips were cleaned in a 5 wt% nitric acid (HNO3) solution followed by ultrasonically rinsing in acetone for 10 min. The acetone was then lightly rinsed from the substrates using Milli-Q® water before being allowed to air dry. After drying the weight of each substrate was recorded using an Ohaus PA214C microbalance.
2.3. Substrate Surface Treatment
The electrolyte used to form the DCPD coatings was prepared by adding 0.32 M of Ca(NO3)2 and 0.19 M of KH2PO4 to 100 mL of Milli-Q® water. The mixture was then thoroughly stirred at 400 rpm for 10 min. The electrolyte was prepared at 25 ± 1ºC and the resulting pH was 4. DCPD coating were prepared by immersing in the prepared electrolyte. Individual substrates were removed from the electrolyte at pre-determined time intervals (3, 15, 30, 60, and 180) so that the size and morphology of the forming coating could be quantified using advanced characterisation techniques. At the end of each immersion period samples were removed from the electrolytic solution, washed in Milli-Q® water, and then allowed to air dry for 24 h. Conversion of DCPD coatings to hydroxyapatite (HAP) was achieved by immersing the DCPD coated substrates into a 1M solution of sodium hydroxide (NaOH) at 80ºC for 2 h. At the end of this period substrates were removed from the electrolyte solution, washed in Milli-Q® water, and then allowed to air dry for at least 24 h. The HAP coatings were then examined using the advanced characterisation techniques discussed below.
2.4. Corrosion Testing
The corrosion resistance of coated and non-coated Mg substrates were tested in freshly prepared phosphate buffer saline (PBS) solution and Ringer’s solution. The composition of the PBS solution (in g/L) consisted of 8.006 NaCl, 0.201 KCl, 1.420 Na2HPO4 and 0.240 KH2PO4. The Ringer’s solution composition (in g/L) consisted of 8.6 NaCl, 0.6 KCl and 0.66 CaCl2. 2H2O. The pH of the respective solutions were adjusted to 7.4 and maintained at 37ºC during the corrosion studies to match body fluid pH and temperature. Polarization curves were generated using an EG&G Princeton Potentiostat/galvanostat (Model 273A, supplied EG&G Princeton applied research) configured for a three-electrode experimental set-up. The working electrode in all corrosion tests consisted of a test substrate with a surface area of 1 cm2. A saturated calomel electrode (SCE) was used as the reference electrode and a platinum wire (Pt) was used as the counter electrode. A Tafel test procedure was performed over a voltage range from -2.5 V up to 1.0 V, with a step size of 10 mV and a 1s time interval for the 10 mV scan rate. From the resulting experimentally derived polarization curves parameters such as corrosion potential (Ecorr), corrosion current density (Icorr), anodic/
cathodic Tafel slopes (βa and βa) and corrosion rate were derived. The polarization resistance (Rp) was calculated at the near open circuit potential (OCP) using the Stern–Geary equation:
𝑅𝑅𝑝𝑝 = (𝛽𝛽𝑎𝑎𝛽𝛽𝑐𝑐)2.303 𝐼𝐼𝑐𝑐𝑐𝑐𝑐𝑐𝑐𝑐 (𝛽𝛽𝑎𝑎+𝛽𝛽𝑐𝑐 )
(4)
2.5. Advanced Characterisation Techniques
2.5.1. X-ray Diffraction (XRD) Spectroscopy
XRD spectroscopy technique was used to examine and identify crystalline size and phases present in the surface coating. Spectroscopy data was recorded at room temperature, using a GBC® eMMA X-ray Powder Diffractometer [Cu Kα = 1.5406 Å radiation source] operating at 35 kV and 28 mA. The diffraction patterns were collected over a 2θ range of 20° to 60° with an incremental step size of 0.02° using flat plane geometry with 2 second acquisition time for each scan.
2.5.2. Scanning Electron Microscopy (SEM) and Energy Dispersive Spectroscopy (EDS)
The SEM technique was used to study the size, shape and morphological features of the surface coating formed on the substrates during the immersion process. All micrographs were taken using a JCM-6000, NeoScopeTM with an attached energy dispersive X-ray spectroscopy function. Samples were mounted on individual substrate holders using carbon adhesive tape before being sputter coated with a 2 nm layer of gold to prevent charge build up using a Cressington 208HR High Resolution Sputter coater.
2.5.3. Transmission Electron Microscopy (TEM) The size, morphology and topography of DCPD particles
formed were investigated using TEM. Sample preparation consisted removing a portion of the coating from the surface of the substrate. The samples were then placed into small tubes containing Milli-Q® water. The tubes were then sealed and placed into an ultrasonic bath for 10 minutes. The suspensions were then filtered 2 times before a single drop from each sample was deposited onto its respective carbon-coated copper TEM grid using a micropipette. The samples were then allowed to slowly dry over a 24-hour period. After sample preparation a bright field TEM study was carried out using a Phillips CM-100 electron microscope (Phillips Corporation Eindhoven, The Netherlands) operating at 80kV.
2.5.4. Fourier Transform Infrared Spectroscopy (FT-IR) FT-IR spectroscopy was used to investigate CaP powders
synthesized during the immersion process using a Perkin–Elmer Frontier FT-IR spectrometer with Universal Single bounce Diamond ATR attachment. FT-IR spectra were recorded in the range from 525 to 4000cm−1 in steps of 4 cm-1.
82 Sridevi Brundavanam et al.: Growth of Flower-Like Brushite Structures on Magnesium Substrates and Their Subsequent Low Temperature Transformation to Hydroxyapatite
3. Results and Discussions 3.1. XRD Analysis of Calcium Phosphate Formation on
Substrates
Before chemical immersion a representative Mg substrate was examined using the XRD technique [49]. The resulting Mg XRD pattern is presented in Figure 1 (c) and confirms the substrate purity. The XRD pattern also revealed that there was no contamination present on the substrate surface. XRD analysis was also carried out on chemically treated substrates. Figure 1 (b) presents a typical XRD pattern of a substrate after an immersion period of 180 min. Analysis of pattern (b) reveals the presence of a crystalline calcium phosphate phase identified as di-calcium phosphate di-hydrate (DCPD) or Brushite (JCPDS 11-293). No other calcium phosphate phases were found in the samples. However, the XRD analysis did confirm the presence of Mg peaks (002), (101) and (102) produced by the underlying substrate. The coating produced during immersion was formed via the following reaction [9, 48]:
HPO42− + Ca2+ + 2H2O → CaHPO4.2H2O (5)
Figure 1. XRD patterns for uncoated and coated substrates: (a) hydroxyapatite coating converted from a DCPD coating; (b) DCPD coating, and (c) pure magnesium substrate
Typically, the mass of DCPD formed during a 180 min immersion period was collected, dried and weighed. The mass was found to be 3.5 mg [Figure 3 (d)] and equated to a coating thickness of around 10 µm. In the next stage of the study, the DCPD coatings were converted into HAP by immersing in a 1M solution of sodium hydroxide at 80 ºC for 2 h. A representative XRD pattern for the converted coating is presented in Figure 1 (a). Analysis of the pattern reveals the DCPD coating was fully converted, and the HAP phases present were consistent with phases listed in the ICDD database. Also present are the Mg peaks seen in the earlier DCPD coatings. The consistent presence of Mg peaks suggests that the coating was unable to fully mask the underlining substrate due to regions of poor surface coverage. The surface coverage was further investigated using SEM analysis and discussed in the following section.
3.2. SEM and EDS Analysis of Coating Formation
The size, morphology and topographical features of the deposited DCPD coatings were examined using SEM and the composition of the coatings were investigated using EDS spectroscopy. Figure 2 presents a series of SEM micrographs taken during the formation of a typical DCPD coating at immersion periods of 3, 15, 30, 60, and 180 min. Figure 2 (a) is a micrograph of a representative Mg substrate prior to chemical immersion. Micrograph 2 (b) was taken after a 3 min immersion period. The image reveals the presence of scattered individual small DCPD plate assemblies that are typically around 50 µm in size. After 15 minutes the scattered plate-like assemblies are beginning to grown into flower-like features. At this point in time the features are around 150 µm in size as seen in Figure 2 (c). For an immersion period of 30 min, the flower-like features have grown in size and are starting to merge with nearby neighbours as seen in Figure 2 (d). Also clearly visible at this stage were large areas of exposed underlining substrate. Following a further 30 minutes of immersion, the flower-like features completely cover the substrate surface as seen in Figures 2 (e) and 3 (a). An enlarged view of the flower-like petals is presented in Figure 3 (b), while Figure 3 (c) presents a detailed end view of the petals showing their plate-like structure. Typically, at the end of a 180 min immersion period substrates were completely covered with large flower-like features typically around 500 to 700 µm in size as seen in Figure 2 (f).
Figure 2. (a) Cleaned magnesium substrate and subsequent formation of DCPD coating after electrolyte immersion periods of (b) 3; (c) 15; (d) 30; (e) 60, and (f) 180 min
The coating formation rate varied over the 180 minute immersion period and is graphically presented in Figure 3 (d). Initially the formation rate was rapid during the first 60 minutes of immersion. Between 60 to 120 minutes the formation rate started slowing down and after 120 minutes
American Journal of Biomedical Engineering 2014, 4(4): 79-87 83
it was significantly reduced as reflected in the levelling off seen in Figure 3 (d). Also seen during the later part of the immersion period was the significant reduction in the evolution of gas bubbles. The significant reduction in formation rate and gas evolution suggests the coating was able to impede the flow of electrolyte towards the substrate surface. Thus, the levelling off seen in Figure 3 (d) after
180 minutes indicates that the coatings were providing some degree of protection to the underlining substrate.
Analysis of DCPD coatings using EDS spectroscopy revealed that they contained elements such as Ca, O and P but no Mg. The analysis indicates that Mg ions were not substituted for Ca ions during the formation of the DCPD coating.
Figure 3. (a) DCPD (brushite) coating decorated with ornate flowers covering the entire substrate surface; (b) enlarged view of flower petals; (c) detailed end view of DCPD crystal plates forming the petal structure, and (d) mass of DCPD coating formed with increasing immersion periods
Figure 4. (a) a representative SEM micrograph of a DCPD flower decorating the coating; (b) EDS analysis of the flower showing its elemental composition; (c) a typical SEM micrograph taken after the conversion to form HAP, and (d) EDS analysis of the coating after conversion
84 Sridevi Brundavanam et al.: Growth of Flower-Like Brushite Structures on Magnesium Substrates and Their Subsequent Low Temperature Transformation to Hydroxyapatite
Moreover, the EDS analysis confirms that the Mg peaks seen in the XRD patterns were the result of the underlining substrate and not Mg in the coating. Figure 4 (a) presents a representative SEM micrograph of a DCPD flower-like structure forming part of the coating, while Figure 4 (b) presents the results of the EDS analysis and the elemental breakdown of components present in the coating.
Figure 4 (c) and (d) present a representative set of results confirming that the conversion process had converted the DCPD coatings into HAP. Figure 4 (d) presents the results of the EDS analysis and the elemental breakdown of the components present in the converted HAP coating. The analysis revealed that calcium to phosphate ratio (Ca:P) ranged from 1.5 to 2.0, with a mean value of around 1.71. The ideal Ca:P ratio for HAP is 1.67 [50], which is approximately 2.4% smaller than the 1.71 determine using the present conversion process.
3.3. FT-IR and TEM Analysis of DCPD Formation
Analysis of the FT-IR spectroscopy data was carried out to identify species, functional groups and vibration modes associated with each peak of the DCPD and HAP coatings. Figure 5 presents the results of the FT-IR analysis. Conversion of DCPD to HAP was confirmed by the appearance of a peak located at 601 cm-1 that corresponds to PO4
3- functional groups normally associated with HAP and not DCPD. TEM analysis of the DCPD coatings revealed that they were composed of plate-like structures that were configured in a flower-like feature. A TEM image of a single DCPD plate taken from a representative coating is presented in Figure 6 (a). Its appearance is solid, angular and typical of DCPD structures seen in the coatings. However, after chemical conversion the appearance and structure of the coating dramatically changes as seen in Figure 6 (b). The solid plate-like structure are transformed into a completely different morphology composed of HAP.
Figure 5. Results of FT-IR analysis showing the presence of PO4
3-
functional groups at peak position 601 cm-1 characteristic of HAP formation
3.4. Corrosion Resistance of Coatings
To be an effective biodegradable material an Mg based
implant must slowly degrade and allow regenerating bone tissues to progressively take over from the load carrying function of the implant. To achieve this objective the corrosion rate of the Mg implant must be effectively controlled and allow sufficient time for successful tissue regeneration to take place. Furthermore, by effectively controlling the corrosion rate of a biodegradable implant means that it is possible to avoid the long-term complications normally associated with conventional metal implants. This study has examined the performance of two CaP coatings in reducing the effects of corrosion on Mg substrates in PBS and Ringer’s solutions at 37ºC. Representative potentio-dynamic polarization curves produced by the corrosion tests for Mg substrates with and without CaP coatings are presented in Figure 7.
Figure 6. TEM images taken before and after conversion: (a) a typical plate-like structure associated with DCPD coating, and (b) a typical HAP structure seen after the conversion process
Figure 7 (a) presents the PBS solution test results and clearly show a significant improvement in corrosion resistance of coated substrates. The DCPD coating provided the superior corrosion resistance compared to the HAP coating. The calculated corrosion rate for DCPD coatings were found to be 0.126 mm/yr and was a significant improvement over the 1.829 mm/yr for the uncoated Mg substrate. Importantly, the corrosion rate of the DCPD coating was half the value found for the equivalent HAP coating as seen in Table 1.
Table 1. Corrosion rates of DCPD and HAP coated Mg substrates determined from polarization curves
Electrolyte Sample Corrosion
Rate (mm/yr)
Humid air [51] Mg substrate 1.0 x 10-5
Distilled water [51] Mg substrate 1.5 x 10-2
Seawater [51] Mg substrate 0.25
Phosphate buffer saline solution (This study)
Mg substrate 1.829
DCPD coated Mg substrate 0.126
HAP coated Mg substrate 0.279
Ringer’s solution (This study)
Mg substrate 3.828
DCPD coated Mg substrate 0.1
HAP coated Mg substrate 0.264
American Journal of Biomedical Engineering 2014, 4(4): 79-87 85
The second electrolyte used in examining corrosion behaviour of substrates was Ringer’s solution. The Ringer’s solution test results are presented in Figure 7 (b). Like the PBS solution results, corrosion testing in Ringer’s solution also confirmed a significant improvement in corrosion resistance of coated substrates. The results also confirmed DCPD coatings were superior to HAP coating. In Ringer’s solution, uncoated Mg substrates had the highest corrosion rate of 3.828 mm/yr. And again DCPD coated substrate’s had a corrosion rate (0.1 mm/yr.) that was slightly less than half of an equivalent HAP coating (0.264 mm/yr.) as seen in Table 1. Corrosion studies confirmed that uncoated substrates found both PBS and Ringer’s solutions highly aggressive. Similarly, body fluids are similar in nature to the test solutions and are as equally aggressive towards uncoated Mg. However, the corrosion studies have shown that both DCPD and HAP coatings have the potential to significantly reduce the degradation rate of Mg substrates in PBS and Ringer’s solutions. It is expected that there would be a similar reduction in corrosion rates for DCPD and HAP coated substrates in vivo. However, further studies are needed to investigate and quantify degradation behaviour in vivo.
Figure 7. Polarization curves produced during the corrosion tests carried out in (a) Phosphate buffer saline (PBS) solution and (b) Ringer’s solution for both DCPD and HAP coated magnesium substrates
4. Conclusions The corrosion rate of Mg substrates in PBS and Ringer’s
solutions at 37ºC was significantly reduced by the presence of a DCPD or HAP coating. The initial DCPD coatings were formed via a straightforward chemical immersion technique. The coatings were subsequently converted to HAP via a low-temperature hydrothermal process in order to examine which coating had the most effective corrosion resistance. FT-IR analysis was used to confirm DCPD conversion to HAP. Corrosion studies revealed that DCPD coated substrates had the lowest corrosion rate in both PBS (0.126 mm/yr) and Ringer’s solution (0.1 mm/yr) compared to HAP coated and uncoated Mg substrates. Both DCPD and HAP coatings have the capacity to significantly reduce the corrosive effects of PBS solution and Ringer’s solution. Since both these solutions are similar in nature to body fluids, the corrosion studies suggest that both coatings have the potential to reduce similar corrosive effects found in the body environment. However, further in vivo studies are needed to fully study and quantify the corrosive effects of actual body fluids on DCPD and HAP coatings formed in this study. This is of particular importance since Mg based degradable implants with DCPD and HAP coatings offer the potential to significantly improve biocompatibility and promote bone formation during tissue regeneration in hard tissue trauma treatments.
ACKNOWLEDGEMENTS The authors would like to thank Mr Ken Seymour for his
assistance with the XRD measurements and Associate Professor Gamini Senanayake for his assistance with the electrochemical measurements.
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55
5.3. Chapter Summary
This chapter examined the degradation behaviour of DCPD and HAP coatings formed
on Mg substrates. Degradation behavioural studies were carried out in PBS and
Ringer’s solutions at 37 ºC. The initial DCPD flower-like coatings were formed via a
straightforward chemical immersion technique. A selection of DCPD coatings was
subsequently converted to HAP coatings via a low-temperature hydrothermal process.
FT-IR analysis was used to confirm the conversion of DCPD coatings to the more
mechanically stable HAP coatings. A range of uncoated Mg substrates, DCPD and HAP
coated substrates underwent degradation studies. The studies revealed DCPD coated
substrates had the lowest corrosion rate in both PBS (0.13 mm/yr) and Ringer’s solution
(0.1 mm/yr) compared to HAP coated substrates [PBS (0.28 mm/yr) and Ringer’s
solution (0.26 mm/yr)] and uncoated Mg substrates [PBS (1.83 mm/yr) and Ringer’s
solution (3.83 mm/yr)]. The research has clearly shown that both DCPD and HAP
coatings have the capacity to reduce the effects of corrosion in both PBS and Ringer’s
solution. The results also suggest that the coatings have the potential to reduce similar
corrosive effects found in the physiological environment. However, this needs to be
further investigated via in vivo studies to fully appraise the corrosive effects on the
respective coatings in the physiological environment. This is of particular importance
since extending the operational life of an Mg based implant is dependent on the
performance of the respective coating in the physiological environment. Furthermore,
any potential enhancement in biocompatibility and bone formation resulting from the
respective coating material needs to be fully investigated in future studies.
56
Chapter 6 – Case Studies 5 and 6: Electrochemical synthesis of
amorphous calcium phosphate coating on magnesium substrates in
aqueous solutions.
6.1. Overview and author contributions
The sixth chapter addresses two components outlined in aims 3 and 4 of the present
research. The first component, listed in aim 3 was to develop and use an electrochemical
technique to produce amorphous calcium phosphate (ACP) coatings on Mg substrates.
In addition, ACP coatings were subsequently converted into HAP coating via
immersion in a sodium hydroxide solution bath at 80 ºC for 2h. The second was to
undertake corrosion studies to investigate the degradation behaviour of the respective
coatings in phosphate buffer saline (PBS) solution and Ringer’s solution at 37 ºC and a
pH of 7.4 to simulate body fluid conditions. Importantly, corrosion studies are needed to
quantify the degradation rate of the coatings, which will determine their suitability in
providing sufficient protection to the underlining substrate during tissue regeneration. In
the ideal case, the coating should slowly degrade during the healing process and at the
same time permit the slow controlled degradation of the underlining Mg substrate. Case
study 5 examines the formation of tube-like structures produced during the
electrochemical process and uses microscopy to assist in developing a formation
mechanism to explain their growth. Case study 6 examines coating formation, coverage
and appraises the degree of protection each coating type was able to provide to the
underlining Mg substrate during the corrosion studies.
S. Brundavanam carried out all electrochemical/corrosion studies and conducted all
advanced characterisation techniques [(XRD), (TEM), (SEM), (EDS) and (FT-IR)] to
57
investigate the structure and surface properties of the ACP and HAP coatings. G.E.J.
Poinern acting as principal supervisor assisted S. Brundavanam in designing the
research program for Chapter 6. S. Brundavanam was also actively involved with the
assistance of D. Fawcett (MANRG - Research Fellow) in analysing the experimental
results and interpreting the advanced characterisation data. S. Brundavanam was
heavily involved in compiling the experimental data and writing the subsequent
research articles. In addition, S. Brundavanam assisted both G.E.J. Poinern and D.
Fawcett with editorial changes recommended by reviewers for the respective peer-
reviewed research articles that form Case Studies 5 and 6. All authors provided
feedback during the preparation of the respective research articles.
6.2. Published Research Articles
Case Study 5
Sridevi Brundavanam, Gérrard Eddy Jai Poinern, Derek Fawcett. Electrochemical
synthesis of micrometre amorphous calcium phosphate tubes and their
transformation to hydroxyapatite tubes. International Journal of Sciences. 2015;
4(2): 7-15.
Case Study 6
Derek Fawcett, Sridevi Brundavanam, Gérrard Eddy Jai Poinern. Growth and
corrosion Behaviour of amorphous micrometre scale calcium phosphate coatings
on magnesium substrates. International Journal of Materials Engineering 2015;
5(1): 10-16.
This article is published under the terms of the Creative Commons Attribution License 4.0
Author(s) retain the copyright of this article. Publication rights with Alkhaer Publications.
Published at: http://www.ijsciences.com/pub/issue/2015-02/
Article Number: v420150206; Online ISSN: 2305-3925; Print ISSN: 2410-4477
Gérrard Eddy Jai Poinern (Correspondence)
+61 8 9360-2892
1Murdoch Applied Nanotechnology Research Group, Department of Physics, Energy Studies and Nanotechnology,
School of Engineering and Energy, Murdoch University, Murdoch, Western Australia 6150, Australia Fax: +61 8
9360-6183
Abstract: Amorphous calcium phosphate tube-like structures were formed on magnesium substrates using a
straightforward electrochemical process. The effect of voltage, time and the resulting hydrogen gas evolution were
found to be responsible for sculpturing the tube-like structures covering the entire substrate surface. Scanning
electron microscopy and Energy Dispersive Spectroscopy was used to quantify the size, shape, structure and
composition of the tube-like structures. Microscopy analysis was used to develop a model to describe the
mechanism behind the formation of the tube-like structures. Also investigated was a hydrothermal technique used to
convert the amorphous calcium phosphate structures into hydroxyapatite. Both X-ray diffraction spectroscopy and
Fourier Transform Infrared spectroscopy were used to confirm the effectiveness of the conversion process.
Keywords: Electrochemical synthesis, amorphous calcium phosphate, microstructures
1. Introduction
Amorphous calcium phosphate (ACP) is the first
mineral phase that precipitates from highly saturated
aqueous solutions containing calcium and phosphate
ions due to its lower surface energy compared to
other calcium phosphates such as octacalcium
phosphate (OCP) and hydroxyapatite (HAP) [1].
Unlike other calcium phosphates, ACP is metastable,
has a non-crystalline character with short-range order
and lacks long-range periodic uniformity normally
associated with crystalline structures [2]. Analysis of
X-ray diffraction measurements have revealed that
ACP has only two very broad and diffuse peaks, with
a maximum occurring at a 2θ angle of 25°. No other
features, normally associated with crystalline
materials are observed [3]. This pattern is typical for
non-crystalline or amorphous materials and confirms
the lack of long-range periodic regularity of ACP.
The spherical structural unit of synthetic ACP is
composed of randomly assembled clusters of ions
approximately 9.5Å in diameter and with the
chemical composition of Ca9 (PO4)6 [4]. Also
contained within the interstices of a typical ACP
particle is tightly bound water that can be as much as
20% by weight of the particle [5]. In aqueous
solutions ACP particles will readily agglomerate, but
are thermodynamically unstable. Therefore, ACP can
be considered a transient or metastable phase that will
easily transform to more thermodynamically stable
crystalline calcium phosphate phases such as HAP
via a process of dissolution, nucleation, and crystal
growth [6]. The amorphous structure of ACP being
unstable is also highly reactive and rapidly
hydrolyses into more stable calcium phosphate
phases in body fluids. It is due to the metastable
properties of ACP that makes it more
osteoconductive than hydroxyapatite and also gives it
superior biodegradability when compared to
tricalcium phosphate when used in vivo [7- 9]. The
existence of ACP in aqueous solutions is governed by
a number of factors such as ionic strength, pH value,
temperature and the presence of contaminants or
additives [10]. Stabilising metastable ACP in aqueous
solutions can be achieved by the introduction of
additive ions or molecules to prevent transformation
to more stable crystalline calcium phosphate phases.
And under in vivo conditions, the presence of protein
based molecules and similar substances can delay the
transformation of ACP due to the effects of kinetic
stabilization [11]. The role of calcium binding
proteins and ionic species involved in the exact
mechanism of biological mineralization and
transformation of ACP to HAP in vivo are not fully
understood [12-14]. However, because of its
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biological compatibility and physiochemical
properties ACP will continue to be used in the rapidly
developing field of tissue engineering, dentistry and
orthopaedics.
This study examines the formation ACP micrometre
scale tubular structures formed on magnesium
substrates during a straightforward electrochemical
synthesis process. The electrodes consist of a
platinum wire mesh (anode) and an Mg substrate
(cathode) immersed in an electrolyte consisting of
calcium nitrate and potassium di-hydrogen
phosphate. Advanced characterisation techniques
such as X-ray diffraction (XRD) spectroscopy,
Transmission electron microscopy (TEM), Fourier
Transform Infrared spectroscopy (FT-IR).
Scanning electron microscopy (SEM) and Energy
Dispersive Spectroscopy (EDS) were used to
determine size, morphology, composition and
architecture of the formed ACP structures.
Furthermore, a proposed formation model is
presented to explain the growth of the ACP
micrometre scale structures.
2. Materials and methods
2.1. Materials
All chemicals used in this work were supplied by
Chem-Supply (Australia) and all aqueous solutions
were made using Milli-Q® water (18.3 M cm-1)
produced by an ultrapure water system (Barnstead
Ultrapure Water System D11931; Thermo Scientific,
Dubuque, IA).
2.2. Mg substrate surface pre-treatment
Magnesium with a purity of 99.9% was used
throughout this study. The sheet was cut into
rectangular strips that were 40 mm in length, 3 mm in
width and 0.15 mm in thickness. The strips were
polishing using 1200 grit silicon carbide (SiC) paper
to remove all surface oxides and contaminants. After
polishing, the strips were cleaned in 5 wt% nitric acid
(HNO3) solution followed by ultrasonically rinsing in
acetone for 10 min. The acetone was then lightly
rinsed from the substrates using Milli-Q® water
before being allowed to air dry. After drying the
weight of each substrate was recorded using an
Ohaus PA214C microbalance.
2.3. Electrochemical formation of ACP surface
features
The formation of ACP surface features on Mg
substrates was accomplished using an
electrochemical cell consisting of two electrodes and
an electrolyte. The electrolyte consisted of an
aqueous solution containing 0.32 M of Ca
(NO3)2.4H2O and 0.19 M of KH2PO4. The electrodes
used in this experimental setup consisted of a
platinum wire mesh (anode) and an Mg substrate that
was used as the cathode. A standard laboratory DC
power supply [GW INSTEK GPS-2303] was used to
supply the required voltage range with a maximum
current density of 60 mA/cm2 in the dynamic mode
setting. Studies were carried out over a voltage range
between 2 and 7 volts and over periods ranging from
30 s to 5 min at a room temperature of 25 ± 1 ºC.
Two voltage/time procedures were followed: 1) fixed
voltage (6 V) and variable time (30 s to 5 min) and,
2) fixed time 3 min and variable voltage (2 to 7 V).
Both procedures were found to give sufficient
experiment data for analysis purposes. After each
experimental run the substrates were removed from
the electrolyte, rinsed using Milli-Q® water and then
air-dried before being stored ready for advanced
characterisation.
2.4. Conversion of ACP structures to hydroxyapatite
Conversion of the ACP structures and coatings to
hydroxyapatite (HAP) was achieved by immersing
the ACP coated substrates into a 1M solution of
sodium hydroxide (NaOH) at 80 ºC for 2 h. At the
end of this period the substrates were removed from
the electrolyte solution, washed in Milli-Q® water,
and then allowed to slowly dry in air for at least 24 h.
Finally, coated Mg substrates were heat treated in an
electric furnace at 250 ºC for 1hour and then allowed
to cool to room temperature in air. The HAP coatings
were then examined using the advanced
characterisation techniques discussed below.
2.5. Advanced characterisation techniques
2.5.1. X-ray diffraction (XRD) spectroscopy
XRD spectroscopy was used to study and to identify
the crystalline phases present in the surface structures
and coating formed during electrochemical synthesis.
Spectroscopy data was recorded at room temperature,
using a GBC® eMMA X-ray Powder Diffractometer
[Cu Kα = 1.5406 Å radiation source] operating at 35
kV and 28 mA. The diffraction patterns were
collected over a 2θ range of 20° to 60° with an
incremental step size of 0.02° using flat plane
geometry with 2 second acquisition time for each
scan.
2.5.2. Transmission electron microscopy (TEM)
The size and morphology of ACP and HAP particles
synthesized were investigated using TEM. Sample
preparation consisted removing a portion of the
synthesized material from the surface of the
magnesium substrate. The material was then placed
into small tubes containing Milli-Q® water. The tubes
were then sealed and placed into an ultrasonic bath
for 10 minutes. The suspensions were then filtered 2
times, and then a single drop from each sample was
deposited onto its respective carbon-coated copper
TEM grid using a micropipette and then allowed to
slowly dry over a 24-hour period. After sample
preparation a bright field TEM study was carried out
Electrochemical Synthesis of Micrometre Amorphous Calcium Phosphate Tubes and their Transformation to
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9
using a Phillips CM-100 electron microscope
(Phillips Corporation Eindhoven, The Netherlands)
operating at 80kV.
2.5.3. Fourier Transform Infrared spectroscopy (FT-
IR)
FT-IR spectroscopy investigations of the synthesized
ACP and HAP were carried out using a Perkin–Elmer
Frontier FT-IR spectrometer with Universal Single
bounce Diamond ATR attachment. Spectra were
recorded in the range from 525 to 4000cm−1 in steps
of 4 cm-1.
2.5.4. Scanning electron microscopy (SEM) and
Energy Dispersive Spectroscopy (EDS)
SEM was used to study the size and morphological
features formed on the substrates during the creation
of the surface coating. All micrographs were taken
using a JCM-6000, NeoScopeTM with attached energy
dispersive X-ray spectroscopy. Samples were
mounted on individual substrate holders using carbon
adhesive tape before being sputter coated with a 2 nm
layer of gold to prevent charge build up using a
Cressington 208HR High Resolution Sputter coater.
3. Results and Discussions
3.1. XRD spectroscopy identification of calcium
phosphate phases
XRD spectroscopy was used to identify the calcium
phosphate phases formed on Mg substrates during the
electrochemical synthesis process. Selections of
typical XRD patterns are presented in Figure 1 (a). In
Figure 1 (a) a representative XRD pattern for an Mg
substrate prior to the electrochemical surface
treatment is presented. Inspection of the pattern
reveals three dominant peaks [(002), (101) and (102)]
and two minor peaks [(100) and (110)] associated
with crystalline Mg and one small peak (101)
associated with Mg(OH)2. Apart from Mg(OH)2 no
other surface contaminants were found on the
substrate surfaces. During electrochemical treatment
a calcium phosphate layer formed over the surface of
the substrates. Subsequent XRD investigation
revealed a characteristic pattern of an amorphous
material as seen in Figure 1 (a). Also present in the
XRD pattern are three peaks [(002), (101) and (102)]
that were identified as Mg from the underlining
substrate and Mg present in the ACP structures. No
other calcium phosphate crystalline phases were
detected in the structured surface coatings.
Subsequent conversion of structured ACP surface
coatings to HAP was also confirmed by XRD
analysis. HAP was confirmed by the presence and
identification of spectra peaks with Miller indices
(002), (211) and (202) that were consistent with
phases incorporated in the ICDD (International
Centre for Diffraction Data) databases. Also present
is the HAP pattern was the presence of (002) related
peak indicating that Mg had been detected. The
strong signal suggests that not only was Mg present
in the substrate coating, but the underlining Mg
substrate was also detected. The extensive and large
crack formations in both the coating and substrate
surface are believed to be responsible for the release
of Mg ions and their subsequent incorporation into
the forming surface structures. The incorporation of
Mg ions into the coatings was also found in the EDS
spectra as seen in Figure 5.
3.2. Microscopy analysis of ACP and ACP based
structural features found in coatings
TEM microscopy was used to examine the structure
of the initial ACP coatings found on the substrates.
XRD analysis identified the formation before
conversion as ACP and TEM images confirm the
spherical structural nature normally associated with
ACP particles as seen in Figures 2 (a) and (b). While
the SEM micrograph presented in Figure 2 (c) shows
the highly agglomerated nature of the ACP particles.
The particles range in size from 60 nm to 250 nm,
with a mean particle size of 120 nm. And because
ACP is thermodynamically unstable, it tends to
transform into more thermodynamically stable
crystalline calcium phosphate phases and this can be
seen in Figures 2 (a) and (b) with the presence of
partially formed plate-like features. The partially
formed plate-like features, which are indicated by
blue arrows, were not detected in the XRD patterns of
the initial ACP coatings. Partially formed plate-like
structures were also seen in SEM images of the
tubular formations as seen in Figure 2 (d), but again
were not seen in the XRD spectra.
3.3. Formation of ACP structures
The electrochemical formation of ACP surface
features on Mg substrates was investigated using two
voltage/time procedures. The first used a fixed time
period (3 min) with a variable voltage range and the
second used a variable time period (30 s to 5 min)
with a fixed voltage of 6 V. The first procedure
investigated the effect of voltage in forming ACP
surface features. The results of this study revealed
that there was a voltage dependence effect and it was
responsible in initiating tubular structure formations
above 4 V. Below 4 V, only small numbers of
hydrogen bubbles were produced and almost no
calcium phosphate materials could be seen forming
except for the occasional small outcrops as seen in
Figures 3 (a) and (b). At this stage the surface of the
Mg substrate was relative smooth and didn’t show
any signs of surface cracking or significant corrosion.
However, from 4 V onwards, gas evolution was
clearly visible across the entire surface of the
substrate. With the dissolution of the Mg substrate
deposits of spherical, granular shaped ACP particles
and small tubules could be seen close to small cracks
that had formed in the surface of the substrate as seen
in Figure 3 (c). The surface cracks seen in the
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substrate are the result of hydroxyl ions accumulating
at the solid/liquid interface [15]. The accumulation of
hydroxyl ions causes a localized region of high pH
that subsequently produces the precipitation of ACP
at the substrate surface. As the voltage was increased,
precipitation of ACP increased and two dominant
features could be seen forming. The first was surface
cracking of the coating that could be seen extending
down to the underling Mg substrate. And the second
feature was widespread tube formation across the
surface of the substrate as revealed in Figures 3 (d) to
(f). An interesting feature of the tubes was the
variation in diameters that ranged from 30 µm up to a
maximum of around 100 µm.
During the second voltage/time procedure the voltage
was maintained at 6 V while the time period ranged
from 30 s to 5 min. After 30 s in the electrochemical
cell, tubes could be seen forming over the entire
substrate surface as seen in Figure 4 (a). In addition
to the tubular features was the presence of numerous
ACP plate-like particles. The plate-like features had
mean diameters of around 5 µm and a thickness of
around 1.5 µm and were covering the surface
between the tubules as seen in Figure 4 (b). Also
present was surface cracks, which were seen earlier in
the variable voltage procedure for voltages above 4
V. One minute into the electrochemical procedure
saw a significant increase in hydrogen evolution and
an increase in size and number of tubules being
formed as seen Figure 4 (c). Three minutes into the
procedure saw a dramatic increase in the size and
depth of cracks in the surface coating, which exposed
more of the underling Mg substrate to the electrolyte,
Figure (d). Also noticeable during this period was the
reduced rate in new tubules being formed as indicated
graphically in Figure 4 (f). At this point in time it is
believed that most of the gas being formed was
exiting the coating via the extensive network of large
cracks covering the surface.
Figure 5 presents a SEM micrograph of a
representative tubular structure formed after 3
minutes under constant voltage. EDS analysis was
carried out on two locations along the length of the
tube. Location (a) was taken at a position three
quarters the length of the tube from the coating
surface and location (b) was taken a quarter of the
length from the coating surface. The spectra resulting
from EDS analysis of both locations is presented in
Figure 5. Both spectra are characteristic of ACP, but
of particular interest is the presence of Mg in the
tubular structure. The presence of Mg in the coating
structures confirms the results of the XRD, which
indicated the presence of Mg. Also present in both
spectra was the presence of oxygen, which suggests
that the presence of Mg in the coating structures
could be in the form of Mg(OH)2. The presence of
Mg(OH)2 in tubular structures could explain the
existence of plate-like features seen in Figure 2 (d).
3.4. Tube Formation Model
Because of the tube-like structures present in the
formation of the deposition layer, it is suspected that
the release of hydrogen bubbles at the substrate
surface (cathode) is one of the key factors responsible
for sculpturing the deposition layer. This assumption
is based on the circular geometry and consistent wall
structure of the tube-like structures that appear to
have been built up around an initial gas bubble
residing on the metal/electrolyte interface. During the
electrochemical process, the Mg substrate acting as
cathode, reduces hydrogen and increases local pH.
With the increase in pH, the stability of HPO42-
significantly increases compared to that of H2PO41-
and as a result the concentration of HPO42- increases
until maximum solubility is reached. Simultaneously,
migration and coalescence of hydrogen at the surface
of the cathode results in the formation of gaseous
macroscopic bubbles as graphically presented in
Figure 6 (a). It is believed that if the bubbles are
small enough they will not be displaced by buoyancy.
At this point, ACP begins to coat the substrate
surface not occupied by bubbles and as a result the
coating acts as an insulator that prevents electron
flow from the substrate. However, the cathodic
substrate continues to supply a current via the
bubble/ACP interface that maintains hydrogen
reduction, ion migration and diffusion. ACP
continues being precipitated from solution and
deposited around surface bubbles. At the same time
hydrogen gas continues to migrate into the forming
tube structure feeding the resident bubble. However,
when the bubble gains sufficient volume, buoyancy
forces dislodge the bubble from the substrate surface.
If there is sufficient hydrogen migration a second
entrapped bubble forms and so one until a steady
stream of small gas bubbles leave the tube formation
site. As the bubble forming process continues, the
length of the tube steadily increases with the self-
assembly of the ACP particles forming the tube wall
structure. If electrochemical conditions are suitable
and sustained then ACP deposition and growth of
tube-like structures continue. However, when gas
migrates into the tube bases ceases there is a rapid in
flow of electrolyte resulting in rapid deposition of
ACP within the tube-like structures.
3.5. Conversion of ACP coatings to HAP coatings
ACP coating structures were converted to
hydroxyapatite (HAP) using a process that involved
immersing substrates into sodium hydroxide at 80 ºC
for 2 h followed by a 1 h thermal treatment in a
furnace at 250 ºC. XRD analysis mentioned above
confirmed that HAP was produced by the conversion
process. Further confirmation of the effectiveness of
the conversion process came from analysis of the FT-
Electrochemical Synthesis of Micrometre Amorphous Calcium Phosphate Tubes and their Transformation to
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IR spectroscopy data as presented in Figure 1 (b).
Conversion of ACP to HAP was confirmed by the
appearance of the peak located at 602.9 cm-1, which
is characteristic of PO43- functional groups normally
associated with HAP and not seen in ACP. In
addition, SEM analysis of the converted coatings
revealed that the tube-like structures had been
transformed into urchin-like structures as seen in
Figures 7 (a) and (b). TEM images of particles
making up the coatings prior to conversion were
spherical and ranged in size from 30 µm to 100 µm in
diameter. After conversion the urchin-like structures
were predominantly composed of needle-like and
plate-like HAP particles. The particle width ranged
from 30 nm to 90 nm, width varied between 80 nm
and 200 nm and the length ranged from 450 nm up to
800 nm. Also present were a small number of
granular particles ranged in size from 30 nm to 100
nm. Figure 7 (c) presents an SEM image of HAP
particles found deposited over the surface of the
substrate and Figure 7 (d) presents a representative
TEM image of HAP particles forming in the urchin-
like structures.
4. Conclusion
A straightforward electrochemical technique was
used to synthesize ACP coatings with micrometre
scale tube-like structures on Mg substrates. The size,
structure and composition of the tube-like structures
were investigated and the subsequent results were
used to propose a formation model. The model
explained the formation mechanism and the
importance of hydrogen evolution in creating tube-
like structures. Also investigated was a hydrothermal
process for transforming the ACP coatings into HAP.
Both XRD and FT-IR confirmed the conversion of
ACP into HAP, with the ACP tube-like structures
being transformed into urchin-like features.
Furthermore, the spherical geometry characteristic of
ACP particles was the prevailing morphology of the
particles prior to the conversion process. After
conversion two HAP particle morphologies were
present. The first consisted of a small number of
granular particles and the second consisted of plate-
like particles. The nanometre scale particles
composed the micrometre scale urchin-like features
that dominated the substrate surface. However,
further studies are needed to fully investigate the
physical and chemical properties of coatings and
structures formed in this preliminary study.
Acknowledgements
The authors would like to thank Mr Ken Seymour for
his assistance with the XRD measurements and Mrs
Sridevi Brundavanam would like to acknowledge
Murdoch University for providing a PhD Scholarship
to undertake her PhD studies.
Disclosure
The authors report no conflict of interest in this work.
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14. Wang L, Nancollas GH. Dynamics of biomineralization
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15. Gray-Munro JE & Strong M. The mechanism of
deposition of calcium phosphate coatings from solution onto magnesium alloy AZ31. Journal of Biomedical Materials Research
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Figure 1 (a) Representative XRD patterns of Mg substrates, ACP and HAP coatings and (b) Results of a
comparative FT-IR study confirming the conversion of ACP to HAP
Figure 2 (a) and (b) TEM images of ACP particles, (c) SEM image of agglomerated ACP particles and (d) ACP
tubular structure composed of ACP particles and partially formed plate-like particles.
Electrochemical Synthesis of Micrometre Amorphous Calcium Phosphate Tubes and their Transformation to
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Figure 3 Growth of ACP tubules over a period of 3 min at voltages of (a) 2, (b) 3, (c) 4, (d) 5, (e) 6 and (f) 7 volts.
Electrochemical Synthesis of Micrometre Amorphous Calcium Phosphate Tubes and their Transformation to
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Figure 4 SEM micrographs recording the growth of ACP tubules using the constant voltage procedure (6Volts): (a)
and (b) 30 s; (c) 1 min; (d) 3 min; (e) 5 min, and (f) number of tubular structures formed as a function of deposition
time.
Figure 5 EDS analysis showing the chemical composition of two separated locations along the length of a typical
tube-like structure formed during the electrochemical process.
Electrochemical Synthesis of Micrometre Amorphous Calcium Phosphate Tubes and their Transformation to
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Figure 6 Tubule formation mechanism: (a) a typical gas bubble residing on the metal/electrolyte interface, (b) build
up of deposited ACP around the bubble, (c) bubble stream formation and (d) wall structure formation.
Figure 7 (a) a representative ACP tube-like structure after HAP conversion; (b) enlarged view of urchin-like
formation after conversion; (c) enlarged SEM image of substrate surface showing HAP particles, and (d) TEM
image of HAP particles formed after conversion.
International Journal of Materials Engineering 2015, 5(1): 10-16
DOI: 10.5923/j.ijme.20150501.03
Growth and Corrosion Behaviour of Amorphous
Micrometre Scale Calcium Phosphate Coatings on
Magnesium Substrates
Derek Fawcett, Sridevi Brundavanam, Gérrard Eddy Jai Poinern*
Murdoch Applied Nanotechnology Research Group, Department of Physics, Energy Studies and Nanotechnology, School of Engineering and Energy, Murdoch University, Murdoch, Western Australia, Australia
Abstract Amorphous calcium phosphate (ACP) coatings were formed on magnesium substrates via a straightforward
electrochemical technique in order to improve the corrosion resistance of the substrates. X-ray diffraction spectroscopy and
microscopy techniques were used to investigate the size, morphology, composition and structure of the ACP coatings.
Analysis of the ACP coatings revealed the presence of micrometre scale fissures and tubular structures. Despite the
presence of these features, the coatings were still capable of significantly reducing the corrosion rate in both PBS and
Ringer’s solutions. Ringer’s solution was found to be the most aggressive towards Mg substrates with a corrosion rate of
3.828 mm/yr. However, after electrochemical treatment, the corrosion rate of substrates coated with ACP was reduced to
0.557 mm/yr. The significant improvement in corrosion resistance is a first step in controlling the corrosion rate of
biodegradable Mg substrates for potential use in hard tissue applications.
Keywords Biodegradability, Electrochemical synthesis, Magnesium, Calcium phosphates
1. Introduction
Magnesium (Mg) is a silvery coloured metal that is
currently being investigated as a potential biologically
degradable implant material for bone tissue engineering.
Mg is a lightweight metal with a high strength to weight
ratio that also has a number of advantageous properties such
as good biocompatibility and biodegradable in the aqueous
physiological environment of the body [1]. In terms of
mechanical properties, Mg is similar in nature to bone in
terms of density and elastic modulus. The density of Mg is
around 1.74 g/cm3 at 20 ºC, while bone tissue ranges from
1.8 to 2.1 g/cm3. In terms of elastic modulus, bone tissue
varies from 40 to 57 GPa and pure Mg around 45 GPa [2, 3].
Interestingly, it is the close similarity between the elastic
modules that makes Mg a potential biocompatible material
suitable for reducing stress shielding and preventing bone
resorption. Stress shielding and the resulting bone
resorption are significant problems that can be encountered
with conventional metallic implant materials [4]. In addition,
unlike Mg, which has been shown to be physiologically
beneficial in the body [5, 6], traditional metallic implants
can release toxic metallic ions such as chromium, cobalt
* Corresponding author:
[email protected] (Gérrard Eddy Jai Poinern)
Published online at http://journal.sapub.org/ijme
Copyright © 2015 Scientific & Academic Publishing. All Rights Reserved
and nickel. The release of these toxic ions during biological
corrosion and mechanical wear often leads to an
unfavourable immune response that significantly reduces
the biocompatibility of the implant and often leads to
secondary revision surgery [7, 8].
However, before Mg can be used in the manufacture of
biodegradable orthopaedic implants two fundamental
problems need to be addressed and both stem from the
metal’s rapid corrosion in chloride rich aqueous body fluids
[9]. The first involves the rapid formation of hydrogen gas,
which overwhelms surrounding tissues and results in the
formation of subcutaneous bubbles [10, 11]. The second
problem results from the rapid loss of mechanical strength
and structural support between the Mg implant and the
surrounding bone tissue that impedes the healing process.
Therefore, effectively reducing the corrosion rate would
significantly improve the mechanical and structural
integrity of the implant for longer periods of time and also
reduce the levels of subcutaneous gas formation. One
interesting method of improving the corrosion resistance of
a susceptible implant material is to coat it with a
non-corrosive protective layer that is both biocompatible
and bioactive. Bioactive coatings based on calcium
phosphates have been successfully applied to coat variety of
conventional metal implants in order to improve their
biocompatibility and reduce the release of corrosion
products into the human body [12].
In this study an amorphous calcium phosphate (ACP)
International Journal of Materials Engineering 2015, 5(1): 10-16 11
[Ca9 (PO4)6] coatings were formed on magnesium substrates
during a straightforward electrochemical synthesis process.
The electrochemical cell consisted of two electrodes. The
first electrode was a platinum wire mesh (anode) and the
second was an Mg substrate that acted as the cathode. Both
electrodes were immersed in an electrolyte consisting of
calcium nitrate and potassium di-hydrogen phosphate. ACP
has a non-crystalline character that lacks long-range
periodic uniformity normally associated with crystalline
structures and readily precipitates from highly saturated
aqueous solutions containing calcium and phosphate ions
due to its lower surface energy [13, 14]. It has a spherical
structural unit composed of randomly assembled clusters of
ions approximately 9.5Å in diameter and also contained
within the interstices is tightly bound water, which can be
as much as 20% by weight of the particle [15]. It is the
metastable properties of ACP in body fluids that makes it
more osteoconductive than hydroxyapatite and also gives it
superior biodegradability when compared to tri-calcium
phosphate [16-18]. And because of its metastable properties,
ACP will transform to a more thermodynamically stable
crystalline calcium phosphate phase such as hydroxyapatite
via a process of dissolution, nucleation, and crystal growth
[19]. Advanced characterisation techniques such as X-ray
diffraction (XRD) spectroscopy, Transmission electron
microscopy (TEM), Scanning electron microscopy (SEM)
and Energy Dispersive Spectroscopy (EDS) were used to
determine size, morphology, crystalline structure and
composition of the ACP coatings formed by the
electrochemical process. The corrosion resistance of the
coatings was then evaluated in phosphate buffer saline (PBS)
solution and Ringer’s solution at body temperature (37ºC).
2. Materials and Methods
2.1. Materials
All aqueous solutions were made using Milli-Q® water
(18.3 MΩ cm-1) produced by an ultrapure water system
(Barnstead Ultrapure Water System D11931; Thermo
Scientific, Dubuque, IA). All chemicals and reagents used
in this study were supplied by Chem-Supply (Australia).
2.2. Surface Pre-treatment of Mg Substrates
High purity (99.9%) Mg foil 0.15 mm in thickness was
cut into rectangular strips 40 mm in length and 3 mm in
width. Silicon carbide (Si C) paper (1200 grit) was used to
remove all surface oxides and contaminants. After initial
surface cleaning, the strips were chemically cleaned in 5%
wt. nitric acid (HNO3) solution followed by ultrasonically
rinsing in acetone for 10 min. This was followed by rinsing
with Milli-Q® water to remove acetone and then allowing
the substrates to air dry. After drying the weight of each
substrate was recorded using an Ohaus PA214C
microbalance.
2.3. Electrochemical Synthesis of ACP Coatings
The electrochemical cell was composed of two electrodes
and an electrolyte. The electrolyte consisted of an aqueous
solution containing 0.32 M of Ca (NO3)2.4H2O and 0.19 M
of KH2PO4. The electrodes consisted of a platinum wire
mesh (anode) and an Mg substrate, which was used as the
cathode. A standard laboratory DC power supply [GW
INSTEK GPS-2303] was used to supply the required
voltage range with a maximum current density of
60mA/cm2 in the dynamic mode setting with a fixed voltage
(6 V) and variable immersion times ranging from 30 s to 5
min.
2.4. Advanced Characterisation of Coatings
2.4.1. X-ray Diffraction (XRD) Spectroscopy
XRD spectroscopy was used to identify the presence of
any crystalline phases present in the surface coating formed
during electrochemical synthesis. Spectroscopy data was
recorded at room temperature, using a GBC® eMMA X-ray
Powder Diffractometer [Cu Kα = 1.5406 Å radiation source]
operating at 35 kV and 28 mA. The diffraction patterns
were collected over a 2θ range of 20° to 60° with an
incremental step size of 0.02° using flat plane geometry
with 2 second acquisition time for each scan.
2.4.2. Transmission Electron Microscopy (TEM)
The size and morphology of ACP particles formed during
the synthesis process were investigated using TEM. Sample
material taken from substrates was placed into small tubes
containing Milli-Q® water, sealed and placed into an
ultrasonic bath for 10 minutes. Suspensions were then
filtered 2 times before a single drop from each sample was
placed onto a carbon-coated copper TEM grid using a
micropipette. After a 24-hour period drying period, samples
were investigated using a Phillips CM-100 electron
microscope (Phillips Corporation Eindhoven, The
Netherlands) operating at 80kV.
2.4.3. Scanning Electron Microscopy (SEM) and Energy
Dispersive Spectroscopy (EDS)
The size and morphology of coating features formed on
Mg substrates were examined using SEM, while the
elemental composition of the coatings was investigated
using EDS. All micrographs were taken using a JCM-6000,
NeoScopeTM with attached energy dispersive X-ray
spectroscopy. Samples were mounted on individual
substrate holders using carbon adhesive tape before being
sputter coated with a 2 nm layer of gold to prevent charge
build up using a Cressington 208HR High Resolution
Sputter coater.
2.5. Corrosion Resistance of ACP Coating
The corrosion resistance of pure Mg substrates and ACP
coating substrates was investigated using freshly prepared
12 Derek Fawcett et al.: Growth and Corrosion Behaviour of Amorphous Micrometre
Scale Calcium Phosphate Coatings on Magnesium Substrates
phosphate buffer saline (PBS) solution and Ringer’s
solution. The composition of the PBS solution (in g/L)
consisted of 8.006 NaCl, 0.201 KCl, 1.420 Na2HPO4 and
0.240 KH2PO4. The Ringer’s solution composition (in g/L)
consisted of 8.6 NaCl, 0.6 KCl and 0.66 CaCl2.2H2O. The
pH and temperature of both test solutions was maintained at
7.4 and 37ºC respectively so that body fluid pH and
temperature could be replicated. An EG&G Princeton
Potentiostat/galvanostat (Model 273A supplied EG&G
Princeton applied research) using a three-electrode
arrangement was used to generate polarization curves. The
working electrode was a test substrate with a surface area of
1 cm2. The remaining two electrodes consisted of a
saturated calomel electrode (SCE) that was used as the
reference electrode, while the counter electrode was a
platinum wire (Pt.). A Tafel test procedure was performed
using a voltage range between -2.5 V and 1.0 V, with a step
size of 10 mV and a 1s time interval for the 10 mV scan rate.
From the resulting experimental polarization curves the
corrosion rate (mm/yr.) was calculated.
3. Results and Discussions
3.1. XRD Spectroscopy and TEM Analysis
Figure 1. XRD analysis of a representative Mg substrate and ACP
coated substrate, with insert TEM image of synthesized ACP particles
The XRD technique was used to examine Mg substrates
before and after the electrochemical process. Figure 1
presents a representative set of XRD patterns for untreated
and treated substrates. Examination of the pattern for the
untreated substrate reveals the presence of six peaks over
the 2θ range from 20º to 60º. The pattern contains three
major peaks [(002), (101) and (102)] and two minor peaks
[(100) and (110)] that correlate to the crystalline structure
of the Mg substrate. Also present in the pattern was a minor
peak (101) that indicated the presence of Mg(OH)2 on the
substrate surface. No other surface contaminants were
found on the substrate. Subsequent electrochemical
processing of substrates produced a thin white calcium
phosphate coating over the surface. The coating was
identified as amorphous calcium phosphate due to the
presence of a very large broad and diffuse peak centred on
the 2θ angle of 25°, characteristic of non-crystalline calcium
phosphate materials [20]. However, three peaks could be
seen in the XRD pattern [(002), (101) and (102)] and were
identified as the underlining Mg substrate.
The inserted TEM image presented in Figure 1 is a
representative of the spherical morphology and highly
agglomerated nature of the ACP particles that make up the
surface coatings of the treated substrates. Particle size
analysis revealed that the particles ranged in size from 60
nm up to 250 nm, with a mean particle size of 120 nm.
3.2. SEM and EDS Analysis of Coating Growth
The SEM technique was used to study the growth of the
ACP coating thickness during the electrochemical process.
The technique was also used to investigate the surface
features and structure of the coatings. Figures 2 (b), (c) and
(d) present micrographs showing the growth and increasing
coating thickness with treatment time. Figure 2 (a)
graphically presents the growth in coating thickness and
reveals that the growth rate varies during the treatment. At
the end of the 5 minute treatment period, a typical coating
reached on average a thickness of around 70 µm. Analysis
of SEM micrographs of the coatings reveals there granular
structure which can be seen throughout the cross-sections.
In addition, with treatment time the coating tends to detach
from the underlining substrate. Also noticeable were the
very rough surfaces of the respective coatings that are
exposed to the electrolyte and the presence of tubular
surface features.
Figure 2. (a) Growth of ACP coating layer; SEM micrographs of mean
thickness measurements taken during growth of coating thickness with
time at 6 V potential (b) 1 min; (c) 3 min, and (d) 70 min
Further SEM investigation of the coating surface
revealed a landscape covered by tubular structures and large
fissures. Figure 3 (a) presents a representative landscape
International Journal of Materials Engineering 2015, 5(1): 10-16 13
showing numerous micrometre scale tubular structures
formed during the electrochemical treatment. Also seen in
Figure 3 (a) are the surface fissures which can be as large as
30 µm in width and zigzag across the surface for many
100’s of µm’s. Many of these fissures, like the tubular
structures provide the electrolyte access to the underling Mg
substrate. Both the fissures and tubular structures have been
formed as a result of hydrogen gas evolution occurring at
the metal-coating interface. Figure 3 (b) presents an
enlarged view of a representative tubular formation seen in
the coating. The central channel diameter of the tube was
estimated to be around 45 µm and extended down through
the coating to the underlining Mg substrate.
Figure 3 (c) presents the results of the EDS analysis and
gives a breakdown of the elemental components present in
the ACP coating. The analysis reveals the presence of
significant amounts of Ca, P and K and traces of O and Mg.
The mean calcium to phosphate ionic ratio (Ca:P) of the
coatings was estimated to be 2.1, which is within the 1.2 to
2.2 range specified by Dorozhkin [13] for ACP. Initially,
the production of ACP was fairly rapid, however, after 3
minutes the deposition rate significantly reduced as seen in
Figure 3 (d). The build-up of ACP after 3 minutes, which
equated to 7.6 mg in mass and a thickness of 55 µm was
significantly reducing the amount of electrolyte in contact
with the underlining substrate and was providing some
degree of isolation from the electrolyte. And by the end of
the 5 minute treatment period the mass of a representative
coating was around 10 mg with a typical thickness of 70 µm.
During the 3 to 5 minute treatment period the coating
formation rate was around 45% lower than the 0 to 3 minute
period. The EDS analysis also found the presence of small
amounts of Mg in the coatings that indicate the possibility
of Mg2+ ions being substituted for Ca2+ ions during the
formation of the coatings. The small amounts of Mg seen in
the EDS analysis may have contributed to, but was not fully
responsible for the intense Mg peaks seen in the XRD.
These strong Mg peaks were primarily the result of the
underlining Mg substrate.
3.3. Corrosion Resistance
The corrosion behaviour of pure Mg substrates and ACP
coated Mg substrates were studied by immersing substrates
into two types of aqueous mediums and then generate
polarization curves using a three-electrode based a
Potentiostat/galvanostat device. Both the PBS solution and
the Ringers solution were maintained at 37 ºC and a pH of
7.4 during the test procedure. Representative polarization
curves produced during the corrosion tests for pure Mg
substrates and ACP coated substrates are presented in
Figure 4. Inspection of the polarization curves of both PBS
and Ringer solutions reveals that both solutions are highly
corrosive to uncoated substrates. For example in seawater,
the corrosion rate is typically around 0.25 mm/yr. [21], the
PBS solution was calculated from the Potentiostat /
galvanostat device to be 1.829 mm/yr. and Ringers solution
was calculated to be 3.828 mm/yr. Clearly from these
measurements it is clear that uncoated Mg substrates will
rapidly corrode in these fluids without protection and
explains why previous attempts of using Mg as an implant
have failed to deliver satisfactory clinical outcomes.
Ringer’s solution was found to be the most aggressive
environment for Mg substrates and this was still the case for
the ACP coated substrates. The ACP coating used in the
corrosion studies was able to reduce the corrosion rate from
3.828 mm/yr. down to 0.557 mm/yr. A similar trend was
also seen for the tests carried out in PBS solution. In this
case the original corrosion rate of 1.829 mm/yr. was
reduced to 0.142 mm/yr. as listed in Table 1.
Figure 3. (a) SEM micrograph of a typical coating surface landscape; (b) enlarged view of a surface tubular structure; (c) EDS elemental analysis, and
(d) deposition rate of ACP coating with time
14 Derek Fawcett et al.: Growth and Corrosion Behaviour of Amorphous Micrometre
Scale Calcium Phosphate Coatings on Magnesium Substrates
Table 1. Corrosion rates of uncoated and ACP coated Mg substrates determined from polarization curves produced from corrosion studies carried out in PBS and Ringer’s solution
Electrolyte Sample Corrosion Rate (mm/yr)
Humid air [21] Mg substrate 1.0 x 10-5
Distilled water [21] Mg substrate 1.5 x 10-2
Seawater [21] Mg substrate 0.25
Phosphate buffer saline solution
(This study)
Mg substrate 1.829
ACP coated Mg substrate 0.142
Ringer’s solution
(This study)
Mg substrate 3.828
ACP coated Mg substrate 0.557
Figure 4. Polarization curves generated during corrosion testing of ACP
coated Mg substrates (70 µm coating thickness) in aqueous environments
at 37 ºC and a pH of 7.4: (a) PBS solution and (b) Ringers solution
The results of the corrosion studies are interesting when
you consider the structure of the ACP coating. There is
clear evidence that the coating integrity is far from
satisfactory, with many micrometre scale surface fissures
and tubular structures present as seen in Figure 3 (a). Both
of which clearly provide access to the underlining Mg
substrate for the respective corrosive solutions. Also seen in
SEM micrographs of coating cross-sections is clear
evidence that the coating is not securely anchored to the
substrate’s surface. Figures 2 (c) and (d) clearly show
significant detachment of the coating, with the coating
lifting as much as 10 µm from the substrate surface.
The most important property of a biodegradable medical
implant material is that it must slowly degrade and allow
regenerating bone tissues to progressively take over from
the load carrying function of the implant. In the case of Mg,
the corrosion rate must be effectively controlled and thus
allowing sufficient time for successful tissue regeneration to
take place.
However, the rapid corrosion rates for uncoated Mg
substrates recorded in this study clearly demonstrate that the
mechanical integrity needed during the healing process
could not be sustained for any significant length of time.
Moreover, the rapid corrosion would produce the evolution
of significant amounts of hydrogen gas, which surrounding
tissues would find difficult to deal with. In spite of this the
corrosion studies of the coated Mg substrates has shown a
significant reduction in the corrosion rates for both PBS and
Ringer’s solutions. The studies have shown that the ACP
coatings formed on the substrate surfaces during the
electrochemical treatment have been able to provide some
degree of protection from aggressive aqueous environments
and reduce the corrosion rate. This result is important since
controlling the corrosion rate of biodegradable Mg
substrates is the first step in developing Mg based implants
for hard tissue applications. It also provides the opportunity
to develop biodegradable Mg based implants that have the
potential to circumvent the long-term complications
normally associated with conventional metal implants
currently used in clinical applications. However, further
research is needed to improve the quality and mechanical
stability of the ACP coatings. I addition, further studies are
needed to study the behaviour of the ACP coating in vivo
since ACP is a metastable calcium phosphate phase. And as
such will easily transform into more thermodynamically
stable crystalline calcium phosphate phases such as
hydroxyapatite via a process of dissolution, nucleation, and
crystal growth [22].
4. Conclusions
A straightforward electrochemical process was used to
form ACP coatings on Mg substrates with the objective of
improving the corrosion resistance of the substrate. XRD
spectroscopy confirmed the formation ACP on the
International Journal of Materials Engineering 2015, 5(1): 10-16 15
substrates while TEM image analysis revealed the
characteristic spherical morphology normally associated
with ACP particles. The image analysis also confirmed the
highly agglomerative nature of the particles. SEM
micrographs taken during the synthesis process reveal the
formation and subsequent growth of micrometre scale
tubular structures and surface fissures in the respective
coatings. Despite the presence of both features, the coatings
were still capable of significantly reducing the corrosion
rate in both PBS and Ringer’s solutions. Ringer’s solution
was found to be the most aggressive with a substrate
corrosion rate of 3.828 mm/yr. However, the substrate
corrosion rate was significantly reduced (0.557 mm/yr.) by
the presence of a 70 µm ACP coating. The significant
improvement in corrosion resistance of ACP coated
substrates is an important result and the present study
demonstrates the viability of using this electrochemical
process to control the corrosion rate. Importantly, this
improvement was achieved in spite of numerous tubular
structures and surface fissures being produced during
coating formation. The disadvantage of the ACP coating
technique is unquestionably the presence of surface fissures
and tubular structures. Conversely, the coatings advantages
include biocompatibility, osteoconductivity, resorption
potential and corrosion resistance. However, further studies
are needed to build on these advantages and perhaps
modifying the current electrochemical process to reduce the
coating imperfections. Thus, being able to control the
corrosion rate of ACP coated Mg substrates is an important
first step in developing a programmable biodegradable Mg
based implant for hard tissue applications.
ACKNOWLEDGEMENTS
The authors would like to thank Mr Ken Seymour for his
assistance with the XRD measurements and Associate
Professor Gamini Senanayake for his assistance with the
electrochemical measurements. Mrs Sridevi Brundavanam
would like to acknowledge Murdoch University for
providing a PhD Scholarship to undertake her PhD studies.
REFERENCES
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[3] Feng, A., and Han, Y., 2010, The microstructure, mechanical
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[4] Ramakrishna, S., Ramalingam, M., Sampath, T. S., Soboyejo, W. O., 2010, Biomaterials: A Nano Approach, Boca Raton, USA: CRC Press, 2010, 161-186.
[5] Saris, N. L., Mervaala, E., Karppanen, H., Khawaja, J. A., Lewenstam, A., 2000, Magnesium. An update on physiological, clinical and analytical aspects, Clinica. Chimica. Acta., 294, 1-26.
[6] Kim, S. R., Lee, J. H., Kim, Y. T., et al., 2003, Synthesis of Si, Mg substituted hydroxyapatites and their sintering behaviours, Biomaterials, 24(8), 1389-1398.
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[9] Mark, P. S., Alexis, M. P., Jerawala, H., Goerge, D., 2006, Magnesium and its alloys as orthopaedic biomaterials: A review, Biomaterials, 27(9), 1728-1734.
[10] McCord, C. P., 1942, Chemical gas gangrene from metallic magnesium, Industrial Medicine, 11, 71-79.
[11] Wen, C. E., Mabuchi, M., Yamada, Y., Shimojima, K. Y., Chino, Y., Asahina, T., 2001, Processing of biocompatible porous Ti and Mg. Scripta Materialia, 45(10), 1147-1153.
[12] Geetha, M., Singh, A. K., Asokamani, R., Gogia, A. K., 2009, Ti based biomaterials, the ultimate choice for orthopaedic implants – A review, Progress in Materials Science, 54, 397–425.
[13] Dorozhkin, S. V., 2009, Nano-dimensional and nano-crystalline apatite’s and other calcium orthophosphates in biological engineering, biology and medicine. Materials, 2, 1975-2045.
[14] Zhao, J., Liu, Y., Sun, W. B., Zhang, H., 2011, Amorphous calcium phosphate and its application in dentistry, Chemistry Central Journal, 5(40), 1-7.
[15] Betts, F., Blumenthal, N. C., Posner, A.S., Becker, G. L., Lehninger, A. L., 1975, Atomic structure of intracellular amorphous calcium phosphate deposits. Proc. Natl. Acad. Sci., 72, 2088-2090.
[16] Li, Y. B., Li, D. X., Weng, W. J., 2007, Amorphous calcium phosphates and its biomedical application, J. Inorgan. Mater., 22, 775-782.
[17] Poinern, G. E. J., Brundavanam, R., Le, X., Nicholls, P. K., Cake, M. A., Fawcett, D., 2014, The synthesis, characterisation and in vivo study of a bioceramic for potential tissue regeneration applications, Scientific Reports, 4 (6235), 1-9.
[18] Dorozhkin, S. V., 2010, Calcium Orthophosphates as Bioceramics: State of the Art, J. Funct. Biomater., 1, 22-107.
[19] Eanes, E. D., Termine, J. D., Nylen, M. U., 1973, An
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Scale Calcium Phosphate Coatings on Magnesium Substrates
electron microscopic study of the formation of amorphous calcium phosphate and its transformation to crystalline apatite. Calcif. Tissue Res., 12, 143-158.
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58
6.3. Chapter Summary
The chapter consisted of two cases studies (5 and 6) that examined the formation of
ACP coating on Mg substrates via an electrochemical technique. The size, structure and
composition of the respective coatings and tube-like surface structures were studied.
Analysis of the data was used to propose a possible tube formation mechanism. The
proposed mechanism explained the importance of hydrogen evolution in forming the
tubes during synthesis. The research also discussed the use of a low temperature
hydrothermal process for transforming ACP coating into HAP. Furthermore, the
degradation behaviour of pure Mg substrates, ACP and HAP coated substrates were
studied in aqueous mediums using a Potentiostat/galvanostat device. The aqueous
solutions used were PBS solution and the Ringers solution. The solutions were
maintained at 37 ºC and a pH of 7.4 during the test procedure. ACP coatings were found
to reduce the corrosion rate from 3.83 mm/yr. (pure Mg substrate) down to 0.56 mm/yr.
in Ringers solution. And a similar trend was also seen in PBS solutions that reduced
significantly the original corrosion rate of 1.83 mm/yr. (pure Mg substrate) down to
0.14 mm/yr.
59
Chapter 7 General discussions, conclusions and future work
Current metallic biomaterials such as stainless steel, cobalt chromium and titanium
alloys are regularly used to replace damaged or diseased bone tissues in clinical load-
bearing applications [1]. At the same time, biologically compatible polymers and
ceramics with lower mechanical strength and lower fracture toughness are routinely
used in low load-bearing applications [2, 3]. Despite the many benefits provided by
each of these biomaterials there still exists problems such as unfavourable inflammatory
responses, non-uniform tissue integration, unfavourable degradation rates and achieving
the appropriate mechanical compliance for long term use. In particular, the release of
detrimental metallic ions and the subsequent adverse inflammatory response can
significantly reduce the biocompatibility of these materials. Also problems related to
stress shielding and bone resorption can also lead to implant failure [4]. And if a
metallic implant is needed for short-term temporary support it often needs to be
removed by a second surgical procedure that significantly increases health costs. The
above mentioned clinical issues clearly demonstrate the need for new biologically
compatible materials that are on one hand capable of providing short-term mechanical
compliance and structural support during the healing process. And on the other hand are
capable of safely degrading with time and hence avoid the issues normally associated
with current long term in situ metallic implants. Research results of the present work
have successfully shown that calcium phosphate (ACP, DCPD and HAP) coatings can
prolong the life of Mg substrates in PBS and Ringers solutions. The results also suggest
that calcium phosphate coated Mg substrates have the potential to be used as
biodegradable implants. Section 7.1 discusses the research outcomes and achievements
60
attained during this thesis. The chapter concludes with recommendations for future
work.
7.1. Research results and achievements
Research undertaken in the present work was designed to develop viable biodegradable
and biocompatible calcium phosphate coatings capable of extending the operational life
of Mg substrates. This thesis is at the forefront of developing a practical and effective
biodegradable Mg implants for potential bone tissue regeneration and bone tissue
engineering applications. The research objectives of the thesis were stated in the four
aims presented in section 1.3 and the subsequent research outcomes were presented as
series of Case Studies.
Chapter 2 presented a comprehensive literature review composed of two articles that
present the current state of research into using Mg based materials for potential
biomedical applications. The first article presented a wide-ranging review of material
properties, biological corrosion mechanisms, alloying elements, surface modifications
and treatments to control the biodegradability of Mg substrates. The second review
article focused on using chemical immersion techniques to produce calcium phosphate
coatings with the potential to reduce biological degradation of Mg based substrates. The
article was of particular importance since it demonstrated how a straightforward wet
chemical based technique could be used to form calcium phosphate coatings on Mg
substrates. The article also presented some preliminary results that showed a significant
improvement in corrosion resistance of DCPD coated Mg substrates when immersed in
a PBS solution with a pH of 7.4 and a solution temperature of 37 ºC.
61
Chapter 3 focused on chemically synthesizing HAP in sufficient quantities to undertake
a number of metal ion adsorption studies. Therefore, Aim 1 of the present work was to
optimise the existing MANRG procedure to produce the necessary quantities of HAP
needed for adsorption studies. Also presented in Chapter 3 was a new synthesis
technique for producing nanometre scale hydroxyapatite powders directly from calcium
hydroxide and phosphoric acid. The new process (Case Study 1) was able to produce
particle sizes ranging from 12.2 nm to 15.4 nm depending on the ultrasonic power used.
The synthesis process was capable of producing significant quantities of HAP and has
the potential to be scaled up to produce larger quantities.
The main thrust of Chapter 4 was to address the objective of Aim 2, which was to
investigate the chemical adsorption capabilities of a highly desirable biocompatible
HAP coating material. The results of these studies form the basis of Case Studies 2 & 3.
The adsorption studies were undertaken to investigate how effective HAP’s complex
hexagonal structure could accumulate metallic ions like the mineral phase found in bone
tissue. The adsorption studies were designed to model the adsorption behaviour of HAP
towards a number of metal ions (Cd2+
, Fe2+
, Cu2+
and Zn2+
) from aqueous solutions and
give an indication of metal ion adsorption capacity of a typical HAP coating. This is of
particular importance because the coating must permit the transfer of metal ions during
substrate degradation. In the case of Mg, besides being biologically degradable, it is
also biologically absorbable and its corrosion products are considered physiologically
beneficial. For example, the body uses Mg in a number of metabolic processes and to
form apatite in the bone matrix [5, 6]. Other beneficial metal ions, in small quantities
include Cu2+
, Fe2+
and Zn2+
. However, not all metal ions are beneficial. For example,
Cd2+
ions readily exchange with Ca2+
ions in bone and results in a number of bone
62
tissue related diseases as discussed in Case Study 2. The results discussed in Case Study
2 revealed that HAP was indeed an effective adsorbent for the removal of Cd2+
ions
from aqueous solutions. The sorption process was endothermic and spontaneous in
nature and followed pseudo-second order kinetics and increased with temperature.
Importantly, the study identified intra-particle diffusion and ion-exchange (Cd2+
→
Ca2+
) as the major participants in the sorption process. While the maximum adsorption
capacity for Cd2+
ions was found to be 123.45 mg/g. Case Study 3 also confirmed intra-
particle diffusion and ion-exchange (metal ion → Ca2+
) were major participants in the
sorption process for Cu2+
, Fe2+
and Zn2+
ions. Fe2+
and Cu2+
ions closely followed
pseudo-second order kinetics and Zn2+
ions tended to follow intra-particle kinetic
diffusion. The maximum adsorption capacities were found to be 61.35 mg/g for Cu2+
ions, 55.25 mg/g for Fe2+
ions and 48.54 mg/g for Zn2+
ions. Both case studies have
confirmed that HAP’s complex hexagonal structure tends to behave in a similar way to
the mineral component found in bone. The results suggest that when HAP is used as a
coating for Mg substrates it will tend to behave in a similar manner as the mineral
component found in the bone matrix. However, future in vivo studies in the appropriate
physiological environment would need to be carried out to fully quantity coating
adsorption, diffusion and desorption processes in situ. This aspect of future research
will be discussed at length in section 7.2.
Chapter 5 presented Case Study 4, which discussed a straightforward chemical
immersion technique and subsequent low temperature thermal treatment to produce
calcium phosphate coatings on Mg substrates. The study also examined the corrosion
resistance of the respective coatings (DCPD and HAP) in PBS and Ringer’s solutions.
The case study addresses two components detailed in aims 3 and 4 of the present work.
63
The first component, listed in aim 3 was to develop and use a chemical immersion
technique to produce calcium phosphate coatings on Mg substrates. The second, listed
in aim 4 was to study the degradation behaviour of each coating type in PBS and
Ringer’s solutions. The chemical immersion technique successfully produced DCPD
coatings that significantly improved the corrosion resistance of Mg substrates. The
degradation studies revealed that the corrosion rate was reduced from 1.83 mm/yr. for
pure Mg down to 0.13 mm/yr. for DCPD coated substrates in PBS solutions. Whereas
degradation studies carried out in Ringer’s solution reduced the corrosion rate from 3.83
mm/yr. for pure Mg down to 0.1 mm/yr. for DCPD coated substrates. HAP coatings
produced via the low-temperature hydrothermal process from DCPD were also found to
significantly reduce the corrosion rate. In PBS solutions, the corrosion rate was reduced
from 1.83 mm/yr. for pure Mg down to 0.28 mm/yr. for HAP coated substrates. While
in Ringer’s solution, the corrosion rate was reduced from 3.83 mm/yr. for pure Mg
down to 0.26 mm/yr. for HAP coated substrates. These results have clearly
demonstrated that both DCPD and HAP coatings are capable of significantly slowing
down the degradation rate of Mg. It is also evident from these studies that both DCPD
and HAP coatings have the ability to prolong the survive-ability of Mg substrates.
However, further in vivo studies are needed to fully investigate degradation processes
operating in the physiological environment and their long-term effect on DCPD and
HAP coated Mg implants. This is of particular importance since the Mg implant is not
intended for long-term placement. Its goal is to slowly degrade and be replaced by
regenerating bone tissue as discussed in future research section 7.2.
Chapter 6 investigated an electrochemical technique that was used to produce ACP and
HAP coatings on Mg substrates. The chapter addresses two components outlined in
64
aims 3 and 4 listed in chapter 1. The first component discussed in aim 3 was to develop
new and straightforward techniques to synthesize calcium phosphate coatings. In this
case an electrochemical technique was used to produce ACP coating on Mg substrates.
In addition, the ACP coatings were subsequently converted into HAP coating via
immersion in a sodium hydroxide solution bath at 80 ºC for 2h. The second component
that formed part of aim 4 was to undertake degradation studies to investigate the
corrosion resistance of the ACP and HAP coating in PBS solution and Ringer’s solution
set at 37 ºC and pH 7.4 to simulate body fluid conditions. Each respective component
forms one of the two cases studies. Both case studies examine the various aspects of
ACP formation and deposition on Mg substrates via the electrochemical technique.
Case Study 5 studied the size, structure and composition of the coating and the
interesting tube-like surface structures via TEM, SEM, EDS and XRD. The results of
the advanced characterisation techniques were used to propose a possible formation
mechanism that was responsible for the formation of the tube-like surface structures
during the electrochemical process. Importantly, the proposed formation mechanism
explained the significance of hydrogen evolution in modelling the architecture of the
tube-like structures as seen in Figure 7.1. Case Study 5 also investigated a low
temperature hydrothermal process for transforming ACP coating structures into HAP.
The two-step process consisted of first immersing ACP coated substrates into a 1M
solution of sodium hydroxide at 80 ºC for 2 h. In the second step the substrates were
thermal treated for 1h in a standard furnace at 250 ºC. Subsequent XRD analysis
confirmed the conversion process and also revealed the presence Mg. The strong Mg
signal in the XRD spectrum suggests that not only was Mg present in the coating, but
the underlining Mg substrate was also detected. SEM and EDS studies also revealed the
presence of extensive cracking in both the surface coatings and underlining substrate
65
surfaces. The extensive cracking is believed to be responsible for the release of Mg ions
and their subsequent incorporation into the forming surface structures.
Figure 7.1 (a) a typical gas bubble residing on the metal/electrolyte interface, (b) build-
up of deposited ACP around the bubble, (c) bubble stream formation and (d) wall
structure formation of ACP.
Case study 6 specifically focused on the coating thickness, mass deposition, coverage,
and appraised the degree of protection or corrosion resistance given by ACP coatings to
the underlining Mg substrate. The corrosion/degradation behaviour was studied in
aqueous based mediums using a Potentiostat/galvanostat device. The two aqueous based
medium used during the degradation studies were PBS solution and Ringers solution.
Both solutions were maintained at 37 ºC and a pH of 7.4 during the test procedure.
Despite the presence of numerous tube-like structures covering the Mg substrate, a 70
µm thick ACP layer could reduce the corrosion rate of 3.83 mm/yr. for pure Mg
66
substrates down to 0.56 mm/yr. in Ringers solution. A similar trend was also seen for
tests carried out in PBS solution. In a PBS solution the corrosion rate for a pure Mg
substrate was found to be 1.83 mm/yr. However, the presence of a 70 µm thick ACP
layer was able to reduce the corrosion rate down to 0.14 mm/yr. In spite of the
promising reduction in the corrosion rate provided by the ACP coatings, SEM cross-
sectional images of the coatings revealed regions were the coating was detached from
the underlining substrate. The presence of coating detachment is of concern and
suggests that further studies are needed to determine coating attachment integrity. The
electrochemical process is capable of producing coatings with the potential to increase
corrosion resistance and reduce Mg degradation. Nevertheless, the presence of coating
detachment over the substrate and the existence of tube-like structures, which extend
down to the underlining substrate are of concern. Both Case Study 5 and 6 have shown
that the electrochemical process in its present form was unsuitable for producing
impervious surface coatings. The presence of tube-like structures that give direct access
to the underlining substrate weakens the mechanical integrity of the deposited coating
and makes this technique unsuitable for producing a stable effective chemical barrier
between the substrate and a surrounding fluid environment. Case Study 4 discussed
earlier, using chemical immersion and a subsequent low temperature thermal treatment
produced a superior coating. Both DCPD and HAP coatings both provided an effective
barrier between the underlining Mg substrate and the surrounding fluid environment
during degradation studies. Both coatings also had a much higher degree of surface
attachment integrity than the ACP coatings seen in Case Studies 5 and 6. Further studies
are of course needed, but the results of the present work have clearly indicates that
chemical immersion was capable of producing a superior coating. Therefore, optimising
the chemical immersion technique and the subsequent low temperature thermal
67
treatment developed in Case Study 4 offers the best method of producing a high quality,
controllable and mechanically stable surface coating on Mg substrates. However, we
still need to investigate better techniques to improve the degradation protection of HAP
coated Mg substrate.
7.2 Recommendations for future work
Research undertaken in this thesis was highly successful in developing viable DCPD
and HAP coatings capable of significantly slowing down the degradation rate of Mg
substrates in PBS and Ringer’s solutions. The research has clearly demonstrated that a
chemical immersion technique and subsequent low temperature thermal treatment could
produce structurally stable coatings capable of prolonging the survive-ability of Mg
substrates in both test solutions. Evaluation and subsequent analysis of the various
physical and chemical properties of each respective coating has revealed that they were
mechanically robust, structurally stable and similar in composition to the mineral
component found in bone tissue. Therefore, the next logical step in developing a viable
biodegradable Mg implant using the coatings produced in the present work would be to
conduct a detailed mechanical characterisation to study the adhesion strength between
coatings and the metal substrate and to evaluate the biological compatibility of the
respective coatings via a series of in vitro cell line studies. This is of particular
importance since it would establish the biocompatibility of the coatings towards cell
types commonly found in the physiological environment. Initially, the in vitro studies
would need to focus on cells directly associated with bone tissues such as osteoblasts
and osteoclasts. A second series of in vitro studies would also need to be carried out to
establish the wider biological compatibility of the coatings towards other cell types
68
found in body tissues. This will be very important, since the coatings should be non-
toxic and non-immunogenic towards all cell types.
In due coarse, if the in vitro studies were positive, the next step would be to undertake
in vivo studies in an appropriate animal model. The subsequent in vivo studies would be
able to give definitive data regarding biological compatibility and coated substrate
degradation rates in a physiological environment. The advantage of being able to
balance the degradation rate is that it will allow bone cells to steadily replace the coated
substrate and at the same time promote the formation of bone tissues and blood vessels
necessary for the delivery of nutrients.
Another avenue for future research would be to investigate using the coatings for the
slow release of pharmaceutical products. The adsorptions studies carried out in the
present work have shown that the HAP coatings were an effective absorber of metallic
ions. Reports in the literature reveal that HAPs structural matrix can store and deliver a
variety of pharmaceutical products such as antibiotics, drugs, enzymes, hormones and
steroids. For example, both HAP-antibiotic and HAP-drug based delivery platforms
have been successfully used in slow release of pharmaceutical products [7-9]. In
particular, the slow and sustained release of pharmaceuticals has been shown to be
effective in the treatment of diseases such as osteomyelitis, osteoporosis and osseous
cancer [10]. Therefore, future studies investigating the use of the substrate coating for
the slow and sustained release of pharmaceuticals during the healing process has
therapeutic benefits.
69
An additional route for future research would be to further investigate the
electrochemical method developed in Chapter 6. Despite the technique being unable to
produce a structurally stable and impervious chemical barrier, the technique did produce
coatings that were able to improve the corrosion resistance of the underlying Mg
substrates. At present the electrochemical technique used only a low voltage DC power
supply. However, an AC or pulsed DC power supply could be investigated in the future
and may be able to deliver an impervious and stable coating. Also, it would be
interesting to compare the metal adsorption data of the synthesised powdered nano HAP
with nano HAP deposited on Mg substrate.
The results of the present research have successfully shown that calcium phosphate
(DCPD and HAP) coatings have the capacity to slow down the degradation rate of Mg
substrates. These coating materials have demonstrated that it is possible for a coated Mg
substrate to safely degrade over extended periods of time and hence avoid the issues
normally associated with current long term in situ metallic implants. Importantly, the
coated substrates have the potential to biologically degrade at a controlled rate and
allow regenerating bone tissues to progressively replace the substrate. Therefore, the
next step in developing a viable Mg based biodegradable metallic implant using the
calcium phosphate coatings developed in the present work would be a series of in vitro
studies to establish material compatibility with a wide range of cell lines. This would be
followed by in vivo studies that would establish the biocompatibility and degradation
behaviour of the coated Mg substrates in a suitable animal model. The results of the
future studies could ultimately lead to clinical application of Mg based biodegradable
metallic implants.
70
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