I.
Introduction recent years, fabrication of porous structures for hard tissue I N augmentation by various synthetic routes using autogeneous cell–tissue interaction has become one of the most promising areas in bone tissue engineering.
1,2
As an alternative to autograft/allograft and xenograft, this simple strategy eliminates the common problems of adverse immune response, donor site morbidity, and pathogen disease transmission. The success depends on whether osteoblasts, chondrocytes, and mesenchymal stem cells are able to be seeded onto a scaffold that can then, over a period of time, degrade and resorb as bone tissue structures grow in vitro and/or in vivo .
3
A three-dimensional (3-D) interconnected porous scaffold is an ideal structure for this purpose, which provides a skeletal support for cell proliferation and differentiation culminating in fast bone in-growth.
4,5 The advantages of a porous bioscaffold lie in the engineering and optimization of material compositions and, in particular, porosity necessary for hard tissue formation without osteolysis. The porosity ranges cited in the literature are highly variable 5–8 but it is generally accepted that a macroporous network of 100–150 m m size benefits from good osteointegration. There are also many reports of pores 4 50 m m in size being used successfully in bone graft substitute (BGS) materials.
9–13
Among porous biocompatible materials, hydroxyapatite
(HAP) having compositions similar to bone, tricalcium phosphate (TCP), and various bioglass compositions are well known
J. Am. Ceram. Soc., 92 [S1] S85–S94 (2009)
DOI: 10.1111/j.1551-2916.2008.02784.x
r 2008 The American Ceramic Society
*
, w
School of Engineering, University of Warwick, Coventry CV4 7AL, U.K.
Highly porous network structures of hydroxyapatite, tricalcium phosphates, a bioactive glass as well as their composites have been fabricated using variations of camphene-, glycerol-, and ice-based freeze-casting techniques. The ball-milled slurries containing 10%–60% solid loading were cast at ambient temperature, followed by sublimation at temperatures between 70 1 and 60 1 C. The green body was sintered in air to a maximum temperature of 1100 1 C for 4 h, which produced excellent threedimensionally (3-D) interconnected structures with open pores.
The nature of the pore channels varied from dendritic, columnar, and cellular to mixed geometry, with dense outer shells in some cases, depending on the particular method used. A monotonic increase in porosity with loading was observed with a decrease in the loading volume. Microstructural analysis was used for porous structures to extract information on geometry, porosity, and pore size distribution. This proved particularly useful to assess whether some of the 3-D structures produced by these methods are suitable for tissue engineering applications. Differential thermal analysis–thermogravimetry, scanning electron microscopy, X-ray diffraction, and density were used to characterize precursor powders, slurry, and sintered products.
for their bioactivity, bioresorption, and osteoconduction.
14
These are used for various purposes as dental and bone void fillers, as a skeletal framework, and as an implant coating to enhance osteointegration and promote bone regeneration. Because the way these compositions resorb varies in a physiological environment, it is instead often desirable to use biocomposites that can be tailored to be site specific where the rates of both resorption and tissue growth can be controlled. A number of bioglass compositions as well as mono-, di-, and
TCP, for example, biodegrade more rapidly than HAP alone.
For clinical applications, porous ceramic–ceramic composite constructs will be more suitable for bone replacement where rapid healing is the desired outcome. It is also well documented that a complex ion-exchange mechanism in body fluid is responsible for the precipitation of an HAP layer on scaffold surfaces promoting osteointegration.
15–17
An important recent finding concluded that the by-products from the dissolution of bioresorbable glasses positively affect osteogenesis via gene expression.
18,19
Numerous fabrication techniques are available to produce ceramic porous structures including the use of porogens,
20,21 replication,
22,23 gel casting, and
24 foaming.
25–29
Generally, these methods produce a highly porous (up to 90%), varying degree of open and closed macroporous structures (100–1000 m m) and interconnectivity but these often suffer from poor mechanical properties, 30 coupled with the lack of control in additive removal during pyrolysis. Alternatively, however, to date, several freeze-casting methods have been used to produce similarly interconnected but microporous structures that exert a positive effect on vascularization, when coupled to resorbable scaffolds in particular. Recently, these have gained considerable attention and, in particular, previous in vivo studies have shown that microporosity can enhance bone growth.
31,32
Also, several groups have reported that both the degree of interconnectivity and the pore size are relevant
9,11,13,33 because a minimum pore size is required for cell infiltration into the scaffold. This work is motivated by the need to explore and apply several freeze-casting techniques that are inexpensive and can be adapted for a low to high solid loadings using a wide variety of bioactive materials to produce 3-D porous structures for applications requiring hard tissue replacement. Therefore, the primary objective of the present study is to investigate the relative merits of a number of freeze-casting techniques including camphene, water and glycerol, and ice mold casting in terms of their ability to generate interconnected porous structures using bioglass, HAP, TCP, and their composites. The specific composite compositions are so chosen as to tailor the material resorption with the cell growth. A comparative analysis of the characterization of the fabricated porous structures is presented, along with a critical evaluation of the nature and influence of the binding phase, pore structure, porosity, and solid loading associated with each method.
G. Messing—contributing editor
Manuscript No. 24681. Received May 16, 2008; approved September 19, 2008.
* Member, The American Ceramic Society
Presented at the 10th International Conference on Ceramic Processing Science (ICCPS-
10), Inuyama, Aichi, JAPAN, May 25–28, 2008 w
Author to whom correspondence should be addressed. e-mail: k.k.mallick@warwick.
ac.uk
S85
II.
Experimental Procedure
(1) Materials
As-received commercial purity (99.9%) HAP, Ca
10
(PO
4
)
6
(OH)
2
, and b -TCP, Ca
3
(PO
4
)
2
, powders (Merck, Darmstadt, Germany) with a median particle size ( d
50
) of 3 m m and a specific surface

S86 area of 72 m
2 derived Bioglass s
Journal of the American Ceramic Society—Mallick
/g were used for all experiments. In house meltgrade 45S5 powder with a mean particle size of o
2 m m was also used. Camphene, C
10
H
6
(Sigma-Aldrich,
Gillingham, UK), alkali-free carboxylic acid (Dolapix CE64,
Zschimmer and Schwarz, Lahnstein/Rheine, Germany), and glycerol (Sigma-Aldrich) were used as the freezing vehicle, dispersant/flocculant, and cryoprotectant (to prevent growth of large ice crystals and crystallized voids), respectively. The average molecular mass of Dolapix CE64 is 320 g/mol, and it is a polyelectrolyte ethanolaminic salt of citric acid.
(5 10
4 mbar) for 24 h. Following sublimation and careful demolding, sintering was carried out at 1100 1 C for 4 h at a ramp rate of 10 1 C/min and the samples were cooled to room temperature.
(4) Thermal Analysis
Vol. 92, No. S1
The isothermal behavior of as-synthesized slurries was determined by simultaneous differential thermal analysis and thermogravimetry (DTA-TGA) (STA1500 TA Instruments, West
Sussex, UK). Approximately 20 mg of slurry was placed in a precalibrated platinum pan and heated to 1000
1
C at a ramp rate of 10
1
C/min in flowing air.
(2) Bioglass and Slurry Preparation
The stoichiometric amount of constituents equivalent to the well-known Bioglass s grade 45S5 composition 45SiO
2
24.5Na
2
O 24.5CaO
6P
2
Hench et al .,
34
O
5 in wt%, originally proposed by was melted in the laboratory in a platinum crucible at 1450
1
C at a heating rate of 50
1
C/min for 4 h. The molten glass was quenched in a pool of freeze-cooled deionized water.
The frit was then micronized from a coarse 30 to
B
2 m m.
Preparation of slurries involved mixing all precursor mixtures investigated with a Dolapix concentration of 6 wt% and ball milling for 48 h using alumina balls in a sealed PTFE jar. The slurry was then cast in a precooled (at 0
1
C) brass split mold, evacuated for 20 min to remove any trapped air bubbles, and stored in a vacuum desiccator at 0 1 C for 24 h, followed by further processing according to the freeze-casting method used.
(5) Phase Analysis and Microstructure
The crystalline phase composition of all porous scaffolds treated at various sintering temperatures and times was determined using powder X-ray diffraction (XRD) in the region of 2 y 5 10 1 –
80 1 with a step size of 0.02
1 and a step duration of 0.5 s on a Nifiltered Philips diffractometer (Model PW1710) using Cu K a radiation ( l 5 0.15406 nm). The evolving phases were matched using an automated powder diffraction software package that included both standard ICDD and calculated ICSD diffraction files.
For scanning electron microscopy (SEM), a Philips Cambridge Stereoscan (Cambridge, UK) and Jeol Model 840 (Jeol,
Watchmead, UK) were used to characterize the pore morphology as well as to observe related microstructural features of porous scaffolds.
(3) Scaffold Fabrication by Freeze Casting
Three separate freeze-casting methods were used to fabricate scaffolds of HAP, a bioglass, and composites of the bioglass and
HAP and HAP and TCP. These methods are referred to hereafter as camphene freeze casting (CFC), water and glycerol freeze casting (WGFC), and ice freeze casting (IFC). The generalized steps used for the fabrication involved mixing of ceramic/bioglass/composite powders with an organic monomer and a dispersant with water/water and glycerol/camphene, followed by deaeration and casting in a precooled brass mold. The cast was carefully removed from the mold and, after relevant freeze-drying processing, the green body was fired up to a maximum temperature of 1100 1 C.
(A) Camphene as a Freeze Vehicle (CFC): A 40 g batch weight of slurry was prepared with solid loadings of 20, 30, 40,
50, and 60 wt% containing the bioglass, a mixture of 70:30 ratio of bioglass/HAP, and a 60:40 ratio of HAP/TCP, and camphene was added as the remainder. A temperature of 60
1
C during ball milling was kept constant to reduce the risk of camphene solidifying and to maintain homogeneity. The demolded green body was freeze dried between 20
1 and 70
1
C, sintered in a muffle furnace at ramp rates of 5
1
–10
1
C/min to a maximum temperature of 1100
1
C (
7
2
1
C) dwelling for 2–4 h, and was finally cooled at the same rate to ambient temperature.
(B) Water and Glycerol as a Freeze Vehicle
(WGFC): A 30 g batch weight of a fixed 40 wt% solid
HAP loading was used to produce a ball-milled slurry. The glycerol content was chosen to be 0, 10, 20, 30, and 40 wt% with water as the remainder. The frozen mold was placed in a thermally insulated container (polystyrene box) into which liquid nitrogen was pumped, thus forming the initial green body after casting. The duration of exposure to liquid nitrogen was 30 min.
The castings in the mold still housed in the polystyrene box were freeze dried for sublimation and subsequently sintered at 1100
1
C for 4 h at a ramp rate of 10
1
C/min.
(C) Ice as a Freeze Vehicle (IFC): A 30 g batch weight of slurry was prepared with solid loadings of 20, 30, 40, 50, and
60 wt% containing a mixture of a 60:40 ratio of HAP/TCP with the water vehicle as the remainder. This method is identical to
CFC, except that slurries were poured into the mold and placed in a thermally insulated container. Liquid nitrogen was poured over the mold in a thermally insulated container and left for 30 min. This was followed by filling the container with dry ice, freeze dried at
B
70 1 C while being pumped in vacuum
(6) p
Pore Analysis
Open porosity, bulk density, and porosity of the green body and the as-fired scaffolds following various sintering treatment protocols were determined using the well-known Archimedes method, pyconometry (Model AccuPyc II 1340, Micromeretics,
Dunstable, UK), and mercury proximity (Autopore IV 9500,
Micromeretics). Depending on the density of the biomaterial used, the porosity p of the scaffold was calculated by
¼ 1 r scaffold r solid where r solid
5 3156 kg/m and 3120 kg/m
3
3 for HAP,
35 for b -TCP.
36
2700 kg/m
3 for bioglass,
35
III.
Results and Discussion
(1) Thermal Behavior
Analysis of the DTA-TGA thermograms indicated that camphene melts at 65
1
C and the slurry admixed with dimethyl carbonate and Dolapix is completely burnt off at around 300 1 C.
An endotherm at 590 1 C and an exotherm at 730 1 C related to the softening point/glass transition ( T g
) and formation of a crystal phase ( T c
) via slow conversion from glass to glass–ceramic.
Melting ( T m
) was observed at around 1050 1 C. These wellknown events have been reported in the literature.
37–39
(2) Phase Evolution and Crystallization
XRD patterns of the HAP, TCP, bioglass, and HAP–TCP,
HAP–bioglass composite powders heated at various temperatures showed that both as-received HAP and TCP powders are phase pure and match the standard JCPDS reference patterns (72-1243/
9-432 and 9-169, respectively). XRD analysis for the ceramic
HAP–TCP composite indicated the presence of the participating constituent phases and their relative intensity ratios, even taking the mass absorption coefficient into consideration, and approximated the actual stoichiometric mixture of the composites.
The crystallization behavior of the unfired as-melted bioglass sample was demonstrated by the appearance of an amorphous
‘‘halo,’’ which is due to the absence of short-range atomic order.
However, at higher sintering temperatures of 850 1 –1000 1 C, the

January 2009 glassy phase begins to surface crystallize to a glass–ceramic with a concomitant increase in peak intensity. This corroborates the thermally induced transformation behavior mentioned earlier.
The crystalline phase, identified as combeite—a sodium calcium silicate, Na
2
Ca
2
Si
3
O
9
, matches the standard JCPDS reference pattern (22-1455) and has been reported in previous studies.
40–44
The effect of a crystallized phase on bioactivity of a glass–ceramic depends on several factors: degree of crystallization, kinetics of hydroxyl calcium apatite (HCA), whose formation is an indicator of the bioactivity, and compositions. There are conflicting reports on the effect of crystallization of bioglass
45S5 on bioactivity. For example, it has been shown that full crystallization can transform this glass to an inert material
45 or to have little or no adverse effect on the bioactivity, except for the rate of surface reaction becoming slower only when the degree of crystallization exceeded 60%.
42,46
It is also shown that while the presence of the Na netics, the formation of the HCA layer was not suppressed.
42,46,47
2
Ca
2
Si
3
O
9 phase depressed the ki-
40–
In the present work, as the crystallization is limited to the surface, with the remainder being mainly an amorphous glassy phase, the level of bioactivity is still expected to be adequate for resorption into bone in-growth.
XRD studies for the bioglass and HAP composite powders showed the structure to be primarily a mixture of crystalline
(HAP) and an amorphous phase, which, on sintering at higher temperatures, transformed into glass–ceramic composites. It was found that the presence of HAP in the precursor mixture somewhat suppresses the extent of glassy phase transforming into the glass–ceramic.
(2) Microstructural Development
Freeze Casting of Porous Bioactive Glass and Bioceramics
(A) CFC: The development of microporosity in scaffolds is shown in typical SEM micrographs in Figs. 1(a)–(c),
Figs. 2(a)–(c), and Figs. 3(a)–(c) for the bioglass and bioglass/
HAP composites, respectively. The striking feature of the camp-
S87 hene processing is the formation of 3-D interconnected pore structures where the morphology of pores is, in all cases, dendritic, generated as a consequence and replication of unidirectional camphene sublimation. The microstructure is quite uniform, where the pore geometry and the general architecture remain unchanged even with a higher solid loading of 60%, as shown in Figs. 1(a)–(c). Although the mean pore size measured is in the region of approximately 65 glass foams.
43 m m, some of the pores, as seen here in the SEM micrographs, are near or above 100 by a controlled and uniform viscous flow at 850 vided glass frit with an average particle size of 2
1 m m m.
The regular coralline nature of pores and the struts is facilitated
C of finely dim. No other studies exist in the literature for porous bioglass and bioglass composite scaffolds fabricated using the camphene method, although a related work reported on replicated macroporous bio-
The effect of the sintering temperature on the development of the porous network in bioglass-derived scaffolds is illustrated in
Figs. 2(a)–(c) for 20 wt% solid loading. The conversion of the bioglass sintered for only 2 h at 850 1 –950 1 C to a glass–ceramic occurred via surface crystallization of Na
2
Ca
2
Si
3
O
9 where.
48
At 950
1 reported to influence the bioactivity of the glass–ceramic and bone regeneration.
40,42
At a lower temperature of 850 1 C, the viscous flow is expected to be low and the pore structure is still intact, as shown in Fig. 2(a). The degree of surface crystallization is minimal, confirmed by the XRD trace still being broad and therefore glassy. This crystallization behavior has been reported else-
C, a relatively higher viscosity resulted in the pore walls becoming denser and the interconnection of the network more defined as seen in Fig. 2(b). At 1050
1
C, close to the onset of the glass melting temperature, the porous network collapses due to nominal viscous flow of the glass–ceramic walls, as shown in Fig. 2(c). The depressed cavities due to melting of glass as well as the crystallized layer are seen instead of the preexisting pores. Although not measured, these samples are likely to show above-average mechanical strength due to the absence of any
Fig. 1.
Typical scanning electron micrographs of bioglass scaffolds produced by camphene-based freeze casting and sintered at 850
1
C using solid loading of (a) 20 wt%, (b) 40 wt%, and (c) 60 wt% [bar 5 100 m m].

S88 Journal of the American Ceramic Society—Mallick Vol. 92, No. S1
Fig. 2.
Scanning electron microscopic microstructure showing the influence of sintering temperature for bioglass scaffolds produced by the camphene freeze-casting technique for a solid loading of 20 wt% at 850
1
C (a), 950
1
C (b), and 1050
1
C (c) [bar 5 200 m m].
Fig. 3.
Typical scanning electron micrographs showing the influence of sintering temperature for bioglass/HAP composite scaffolds produced by the camphene freeze-casting technique for a solid loading of 20 wt% at 850
1
C (a), 950
1
C (b), and 1050
1
C (c) [bar 5 100 m m].

January 2009 Freeze Casting of Porous Bioactive Glass and Bioceramics S89
Fig. 4.
Effect of measured porosity with solid loading for scaffolds sintered at 850
1
C ( & , bioglass; , bioglass/HAP composite; m , HAP/TCP composite).
Fig. 5.
Plot of mean pore size of scaffolds as a function of solid loading heat treated at 850
1
C ( & , bioglass/HAP composite; , bioglass; m ,
HAP/TCP composite).
defects, and the compressive strength is estimated to be around
0.3–0.7 MPa.
43,44
Therefore, an optimum sintering temperature of 850 1 –950 1 C for fabricating such glass–ceramic scaffolds is preferable in order to generate an arrayed interconnected porous network structure. Because these bioactive glasses are known to elicit cell proliferation, this should result in better bone in-growth also tailored to the rate of biodegradation, i.e. via controlled biodegradability
35,48,49 of bioglass and related glass– ceramics when implanted. Various reports supporting this are available in the literature.
34,50–53
The effect of the sintering temperature on the HAP/bioglass composite scaffolds is shown in typical SEM micrographs in
Figs. 3(a)–(c). Here, as the composites predominantly contain ceramic HAP (70 wt%) compared with the bioglass (30 wt%), viscosity is restricted due to ceramic particles having a very high melting temperature. At 850 1 C, the porous microstructure is similar to that of the bioglass at this temperature but the pore boundaries appear to be somewhat undefined. The pore structure at 950 1 C has changed significantly, wherein particulates no longer restricting the flow of the glass HAP and the pore channels, produced originally by the directional sublimation of camphene are now more delineated. At 1050
1
C, the composite structure, in contrast to bioglass, does not collapse due to HAP particles well supporting the porous network even though melting of the bioglass is already underway. The thickness of the pore walls varied from 8 to 10 m m, producing networked pores, thought to be formed around the HAP particles via melting.
Thus, fabrication and reproducibility of the network structure would depend on two interlinked factors: precursor particulate size and heat-treatment temperature. For bioglass, it was observed, from a preliminary study during this investigation, that the overall porosity is usually about 20%–30% lower for starting particle size 4 30 m m compared with a size range of 1–4 m m.
This is supported by previous studies on calcium phosphate glass where the difference in porosity can be as high as 10% between large and fine particle sizes of 15–20 and 1–5 m m, respectively.
54
Clearly, an optimum and preferred sintering temperature of up to 1000
1
C can easily be tolerated especially for a lower solid loading and a finely divided size distribution at the expense of perhaps a slight reduction in porosity.
The level of porosity of the bioglass and bioglass/HAP composites was found to be dependent on the sintering temperature and loading as mentioned earlier. Porosity varied linearly with a solid loading of 20%–60%, as shown in Fig. 4. A maximum of about 70% porosity was achieved for the bioglass. However, only 64% porosity could be achieved for the composite samples due to limited viscous flow. An almost uniform strut thickness for the lowest solid loading (20%) was an average of 5 m m, which increased with the loading level as the porosity decreased.
The number of dendritic bridges/struts between the pores increased while pores became smaller with increased solid loading.
The trend in the linear variation of the mean pore size with loading is shown in Fig. 5. At 950 1 C, the mean pore sizes were
50 and 60 m m for the bioglass and its composites, respectively, for 20% loading, but resulted in a much reduced level of microporosity of around 10 m m at 60% solid loading. A possible explanation for a difference in approximately 10% in porosity may lie, apart from the increased glass viscosity, in the variation of cooling rate, i.e. temperature decrease in unit time during sublimation. Although the processing protocol was empirical, it is possible that there may well be a temperature differential in the liquid nitrogen-precooled molds between the two groups. As a result, the molds equilibrated at different rates to the final freeze-drying temperature. Also, the mechanisms of vapor transport and mass transport behavior during sublimation
55 for the bioglass and the composites must have contributed to the different freeze-drying rates due to different, but unavoidable, conditions of cooling. In addition, the mean pore diameter increased with the increase in sintering temperature as the glass started to flow and produce a more delineated highly interconnected 3-D porous structure. This is true, except for bioglass scaffolds, whose network collapses at 1050
1
C.
Figures 6(a)–(d) show typical SEM micrographs of the porous architecture of HAP/TCP composites, at different solid loadings, fabricated at 1100
1
C. The general microstructure seen in Fig. 6(a) can be described typically as cellular type and networked although there is some evidence, as indicated in
Fig. 6(b), of localized channel/lamellar structures, a feature not observed in samples of bioglass and its composites. It is likely that the vapor and/or the mass transport mechanism, during sublimation, is relatively more stable thermodynamically in this mixed composite ceramic phase than is the case for the mixed ceramic–glass composite samples. There is also clear evidence of microporosity on the column walls with variation in pore diameters (20–40 m m) and microporous pits of
B
1–2 m m, as shown in Fig. 6(c). The structure is less porous due to densification but still partially cellular at a loading as high as 60 wt%
Fig. 6(d). For all loadings, the increase in porosity and pore size is monotonic with decreasing loading level and exhibits a trend very similar to that observed for bioglass composites, as shown in Figs. 4 and 5.
(B) WGFC: For HAP freeze casts, the effect of water and glycerol as a chosen solvent vehicle is shown in typical SEM micrographs in Figs. 7(a)–(f). The microstructures are generally very similar to that observed for the camphene technique, where the fabricated structure is highly porous and interconnected. A noticeable difference, however, is the increased regularity and homogeneity of the porous structure showing dendritic or coralline-like geometry and the formation of denser pore walls. This is shown in Figs. 7(a)–(d). The walls were without defects, thus increasing the size of the pores, and are due to water acting as a solidification modifier, which helps eliminate defects associated

S90 Journal of the American Ceramic Society—Mallick Vol. 92, No. S1
Fig. 6.
Typical scanning electron micrographs of the porous architecture of HAP/TCP composite scaffolds produced by camphene-based freeze casting and sintered at 1100
1
C for 4 h at solid loading of (a) 20 wt%, (b) and (c) 30 wt%, and (d) 60 wt% [bar 5 1 mm (a); bar 5 500 m m (b); bar 5 20 m m
(c), (d)].
with the particle rejection during freeze casting (i.e. defects in the pore walls). Also, no voids were observed, although several reports suggested that the particle rejection can result in the formation of large voids.
56,57
The lower cooling rate from liquid nitrogen during casting and solvent sublimation created a reproducible pore structure. It can be seen from Fig. 8 that the relationship of both porosity and pore size is monotonic with increasing glycerol concentration. The dendritic bridges are evident regardless of the solvent concentration. A maximum of
55% porosity was achieved with a wide variation of mean pore diameters between 10 and 50 m m exhibiting a homogenous interconnected wall thickness of around 8 m m, considering the solid loading is as high as 40 wt%. The smaller pore diameter is due to the effect of glycerol, which is known to reduce the size of the ice crystallites.
57–59
The technique demonstrates that a uniform level of microporosity can be controlled by tuning the glycerol/water content in the slurry.
An interesting observation was made for pore structure with an average pore diameter of 45 m m, shown in Fig. 7(e), in one of the runs for 30 wt% glycerol concentration. A magnified image in Fig. 7(f) shows the presence of a large number of interconnected microporous surface pits (2–3 m m) on the dendritic bridges. Such pits have been observed previously 60 for 5%–
10% solid loading of HAP using the camphene method. However, the reported pits were not connected as is the case here. It is suggested that microscopic frozen ice crystals separated by glycerol molecules formed, on sublimation, such protrusions in certain preferred growth directions such as
/
100
S and
/
111
S
.
61
Formation of such micropores is desirable and is expected to accelerate the cellular response.
62
Although the literature on the use of water and glycerol in producing a porous HAP scaffold is scarce, a recent study reported
58 constructs of only 1–2 m m pores and porosity of o
50%. In contrast, it should be noted here that the present study achieved the control of microporosity for a very high bioceramic loading using a hydrogen bond-forming cryoprotectant. The freezing behavior was similar to those used for aqueous slurries of alumina.
57,63
(C) IFC: The effect of ice as a freezing vehicle on the development of pore structure for HAP/TCP composite is shown in SEM micrographs in Figs. 9(a)–(f) and Figs. 10(a)–
(b). Typical examples in Figs. 9(a) and (b) show that macroporous columnar channels formed on sublimation of the advancing solidification/growth front of the ice. A majority of these were observed only in samples up to 30% loading. The 3-D periodic channels were continuous and aligned parallel to the direction of ice growth, although some were also formed in the perpendicular direction. It is likely that these were formed due to a slight variation in the rate of sublimation of the ice dendrites, thereby inducing a degree of directional cooling as the temperature was increased from liquid nitrogen temperature to 70 1 C.
It is also reported that microstructural evolution to planar
(dense), cellular, and lamellar zones is sequential and is directly related to the initial freezing front.
30
This is why it is often difficult to ascertain the transition from one zone to the other as the freezing ice front maintains the morphology. The interchannel/ interlamellar width of the open interconnected macropores typically ranged between 30 and 50 m m, with an average channel wall thickness of 5–10 m m. The surface morphology on the internal macroporous channel walls exhibited a homogeneous and interconnected dendritric relief, as shown in Figs. 9(c)–(d), varying in pore size (5–10 m m), and can be technically regarded as microporous. The observed dendritic branch-like structure, seen in Figs. 9(e) and (f), is a culmination of microscopic ice formation, and the width of the struts between dendrites was typically
3 m m. The dendritic arms also contain micropores (1–2 m m) as pits. Formation of such dendritic structures has also been reported for an ice-templated porous alumina structure.
30,64
It should be noted here that the presence of such a diverse array of microporous structure should be beneficial in both vascularization and proliferation of bone cells.

January 2009 Freeze Casting of Porous Bioactive Glass and Bioceramics S91
Fig. 7.
Scanning electron microscopic microstructure of the water- and glycerol-based freeze cast HAP/TCP composite scaffolds with a solid loading of
40 wt% showing the influence of glycerol concentration in water with (a) without glycerol (b) 10 wt%, (c) 20 wt%, (d) 30 wt%, (e) atypical features at 30 wt% and (f) magnified image of (e) [bar 5 200 m m (a), (c), (e); bar 5 100 m m (b) and (d); bar 5 5 m m (f)].
As the composite particulate concentration increased with solid loading and the content of ice in the frozen body reduced, the pore structure became microporous, interconnected, and homogeneous. In this case, more nucleation sites were available for the ice crystals to grow and fewer water molecules sublimated, leading to a progressive reduction in porosity with loading, as shown in Figs. 10(a) and (b). The pore for 50% loading was a hybrid cellular/dendritic with varying aspect ratios and resembled the dendritic surface morphology as observed in
Fig. 9(a). Such a morphology has been reported for high loading (80%) of an alumina structure using a method similar to the current work.
65
With higher loading, densification occurred, but the structure was still composed of a porous 3-D network.
The porosity is directly dependent on the solid loading, which, in turn, is related to the water content in the slurry.
The observed microstructural features are supported by the variation of total porosity with loading, where the relationship is one of linearity. A maximum of 75% porosity was achieved with the least loading (20%) possible. It was demonstrated that this technique produced a wide range of pore sizes, between 20 and 100 m m, and a mean pore diameter of 60 m m.
Fig. 8.
Influence of solvent concentration on porosity and pore size of
HAP/TCP composite scaffolds at 40 wt% solid loading ( , porosity; & , pore size).

S92 Journal of the American Ceramic Society—Mallick Vol. 92, No. S1
Fig. 9.
Scanning electron microscopic microstructure of HAP/TCP composite scaffolds for solid loading sintered at 1100
1
C for 4 h showing (a) and (b) columnar macrochannel, (c) and (d) channel wall pores, and (e) and (f) microporous dendrites [bar 5 50 m m (a); bar 5 20 m m (b); bar 5 5 m m (c), (d); bar 5 2 m m (e), (f)].
IV.
Conclusions
Using several freeze-casting techniques with various slurry concentrations, the author has successfully fabricated highly porous
3-D network structures of a bioactive glass and its composites with HAP as well as bioresorbable HAP/TCP composites. Although SEM study has established the interconnectivity between pores, the degree of interconnectivity has not been fully determined and future work using m CT is planned to evaluate this aspect. The effectiveness of techniques such as camphene, water and glycerol, and IFC to produce a porous scaffold network has been demonstrated, and the results have been evaluated on the basis of a comparative analysis of the pore structure of the 3-D constructs. It is concluded that all the methods used are suitable with varying degrees of control to fabricate structures, which, among other applications such as BGS, may be adapted for bone tissue engineering. Therefore, the following conclusions can be made:
(1) The study has demonstrated the fabrication of bioactive glass-based 3-D and microporous scaffold structures. Structures with porosity ranging from 50% to 70% have been successfully fabricated using the well-known camphene-based freeze-casting method.
(2) Formation of a defect-free and interconnected porosity of a maximum of 55% was achieved for HAP-based composites using WGFC. At all solvent concentrations, it was possible to replicate a coralline-like uniform dendritic with isolated evidence of micropores of 2–3 m m size on the interconnected bridges. This was found to be facilitated by very low temperatures during the sublimation process.
(3) The IFC technique produced highly porous and networked composite structures, ranging from a columnar to a dendritic architecture. The nature of pore morphology can be understood in terms of a sequential process of initial freezing/ solidification front progressing through various zones dictated essentially by the physics of ice crystals. The porosity achieved was as high as 75% and the level was maintained to nearly 50% relative to a higher solid loading of 60%.
(4) While it is acknowledged that 100–150 m m is an optimum macroporous range for bone in-growth, microporosity

January 2009 on complex physiological factors.
Freeze Casting of Porous Bioactive Glass and Bioceramics
Fig. 10.
Scanning electron microscopic microstructure of HAP/TCP composite scaffolds showing channel morphology formed at solid loading of (a) 50 wt% and (c) 60 wt% [bar 5 50 m m (a), (b)].
may enhance in vivo bioactivity and bone in-growth. Also, based on the literature, neovascularization is expected to be promoted by the presence of the observed microporosity. Some of the structures produced in the present work fall within the microporous range and are just outside the minimum macroporous values. However, it should be kept in mind that the fabricated structures are bioresorbable with varying degrees and as resorption progresses the porosity is also expected to increase depending on the degree and the rate of HCA formation. This in turn is related to the choice of a preferred biomaterial. Therefore, the pore size range should not be regarded as absolute but depends
Acknowledgments
The author thanks Chris Ashurst and Amy Kamara for their help with some of the experimental arrangements and Martin Davis for running samples for XRD, thermal analysis, and SEM.
References
1
C. W. Patrick Jr., A. G. Mikos, and L. V. McIntire, ‘‘Prospectus of Tissue
Engineering’’; pp. 3–14 in Frontiers in Tissue Engineering , Edited by C. W. Patrick
Jr., A. G. Mikos, and L. V. McIntire. Elsevier Science, New York, USA,
1998.
2
K. J. L. Burg, S. Porter, and J. F. Kellam, ‘‘Biomaterial Developments for
Bone Tissue Engineering,’’ Biomaterials , 21 [23] 2347–59 (2000).
3
4
R. Langer and J. P. Vacanti, ‘‘Tissue Engineering,’’ Science , 260 , 920–6 (1993).
L. L. Hench and J. M. Polak, ‘‘Third-Generation Biomedical Materials,’’
Science , 295 [5557] 1014–7 (2002).
5
J. R. Jones and L. L. Hench, ‘‘Regeneration of Trabecular Bone Using Porous
Ceramics,’’ Curr. Opin. Solid State Mater. Sci.
, 7 [4–5] 301–7 (2003).
S93
6
R. E. Holmes, ‘‘Bone Regeneration within a Coralline Hydroxyapatite Implant,’’ Plast. Reconstr. Surg.
, 63 , 626–33 (1979).
7
O. Gauthier, J. M. Bouler, E. Aguado, P. Pilet, and G. Daculsi, ‘‘Macroporous
Biphasic Calcium Phosphate Ceramics: Influence of Macropore Diameter and
Macroporosity Percentages on Bone Ingrowth,’’ Biomaterials , 19 , 133–9 (1998).
8
J. Bobyn, R. Piliar, H. Cameron, and G. Weatherly, ‘‘The Optimal Pore Size for the Fixation of Porous Surfaced Metal Implants by the Ingrowth of Bone,’’
Clin. Orthop. Rel. Res.
, 150 , 263–70 (1980).
9
P. S. Eggli, W. Muller, and R. K. Schenk, ‘‘Porous Hydroxyapatite and Tricalcium Phosphate Cylinders with Two Different Pore Size Ranges Implanted in the Cancellous Bone of Rabbits. A Comparative Histomorphometric and Histologic Study of Bony Ingrowth and Implant Substitution,’’ Clin. Orthop. Rel. Res.
,
232 , 127–38 (1988).
10
K. A. Hing, S. M. Best, K. E. Tanner, W. Bonfield, and P. A. Revell,
‘‘Mediation of Bone Ingrowth in Porous Hydroxyapatite Bone Graft Substitutes,’’
J. Biomed. Mater. Res.
, 68A [1] 187–200 (2004).
11
R. E. Holmes, V. Mooney, R. Buchholz, and A. Tencer, ‘‘A Coralline Hydroxyapatite Bone Graft Substitute,’’ Clin. Orthop. Rel. Res.
, 188 , 252–62 (1984).
12
J. J. Klawitter, J. G. Bagwell, A. M. Weinstein, B. W. Sauer, and J. R. Pruitt,
‘‘An Evaluation of Bone Growth into Porous High Density Polyethylene,’’
J. Biomed. Mater. Res.
, 10 , 311–23 (1976).
13
J. H. Kue`hne, R. Bartl, B. Frish, C. Hanmer, V. Jansson, and M. Zimmer,
‘‘Bone Formation in Coralline Hydroxyapatite. Effects of Pore Size Studied in
Rabbits,’’ Acta Orthop. Scand.
, 65 [3] 246–52 (1994).
14
L. L. Hench, ‘‘Bioceramics: From Concept to Clinic,’’ J. Am. Ceram. Soc.
, 74
[7] 1487–510 (1991).
15
M. Vallet-Regi, A. Ramila, S. Padilla, and B. Munoz, ‘‘Bioactive Glasses as
Accelerators of Apatite Bioactivity,’’ J. Biomed. Mater. Res.
, 66A [3] 580–5 (2003).
16
L. L. Hench, ‘‘Bioactive Materials: The Potential for Tissue Regeneration,’’
J. Biomed. Mater. Res.
, 41 [4] 511–8 (1998).
17
M. M. Pereira, A. E. Clark, and L. L. Hench, ‘‘Calcium Phosphate Formation on Sol–Gel-Derived Bioactive Glasses In Vitro,’’ J. Biomed. Mater. Res.
, 28 [6]
693–8 (1994).
18
I. D. Xynos, A. J. Edgar, L. D. Buttery, L. L. Hench, and J. M. Polak, ‘‘Gene-
Expression Profiling of Human Osteoblasts Following Treatment with the Ionic
Products of Bioglass s
45S5 Dissolution,’’ J. Biomed. Res.
, 55 [2] 151–7 (2001).
19
I. D. Xynos, A. J. Edgar, L. D. K. Buttery, L. L. Hench, and J. M. Polak,
‘‘Ionic Products of Bioactive Glass Dissolution Increase Proliferation of Human
Osteoblasts and Induce Insulin-Like Growth Factor II mRNA Expression and
Protein Synthesis,’’ Biochem. Biophys. Res. Commun.
, 276 , 461–5 (2000).
20
O. Lyckfeldt and J. M. F. Ferreira, ‘‘Processing of Porous Ceramics by Starch
Consolidation,’’ J. Eur. Ceram. Soc.
, 18 , 131–40 (1998).
21
M. Boaro, J. M. Vohs, and R. J. Gorte, ‘‘Synthesis of Highly Porous Yttria-
Stabilized Zirconia by Tape-Casting Methods,’’ J. Am. Ceram. Soc.
, 86 [3] 395–4
(2003).
22
K. Schwartzwalder and A. V. Somers, ‘‘Method of Making Porous Ceramic
Article’’; U.S. Patent No. 3090094 1963.
23
K. Jun, Y. H. Koh, J. H. Song, S. H. Lee, and H. E. Kim, ‘‘Improved Compressive Strength of Reticulated Porous Zirconia Using Carbon Coated Polymeric
Sponge as Novel Template,’’ Mater. Lett.
, 60 [20] 2507–10 (2006).
24
P. Sepulveda and J. G. P. Binner, ‘‘Processing of Cellular Ceramics by Foaming and In Situ Polymerization of Organic Monomers,’’ J. Eur. Ceram. Soc.
, 19 ,
2059–66 (1999).
25
C. Tuck and J. R. G. Evans, ‘‘Porous Ceramics Prepared from Aqueous
Foams,’’ J. Mater. Sci. Lett.
, 18 , 1003–5 (1996).
26
S. Dhara and P. Bhargava, ‘‘A Simple Direct Casting Route to Ceramic
Foams,’’ J. Am. Ceram. Soc.
, 86 , 1645–50 (2003).
27
K. Prabhakaran, N. M. Gokhale, S. C. Sharma, and R. Lal, ‘‘A Novel Process for Low-Density Alumina Foams,’’ J. Am. Ceram. Soc.
, 88 [9] 2600–3 (2005).
28
M. Pradhan and P. Bhargava, ‘‘Effect of Sucrose on Fabrication of Ceramic
Foams from Aqueous Slurries,’’ J. Am. Ceram. Soc.
, 88 [1] 216–8 (2005).
29
T. Fukasawa, M. Ando, T. Ohji, and S. Kanzaki, ‘‘Synthesis of Porous
Ceramics with Complex Pore Structure by Freeze-Dry Processing,’’ J. Am. Ceram.
Soc.
, 84 [1] 230–2 (2001).
30
S. Deville, E. Saitz, and A. Tomsia, ‘‘Freeze Casting of Hydroxyapatite
Scaffolds for Bone Tissue Engineering,’’ Biomaterials , 27 [32] 5480 (2006).
31
K. A. Hing, ‘‘Characterisation of Porous Hydroxyapatite,’’ J. Mater. Sci.:
Mater. Med.
, 10 , 135–45 (1999).
32
K. A. Hing, ‘‘Microporosity Enhances Bioactivity of Synthetic Bone Graft
Substitutes,’’ J. Mater. Sci.: Mater. Med.
, 16 , 467–75 (2005).
33
P. A. Rubin, J. K. Popham, J. R. Bilyk, and J. W. Shore, ‘‘Comparison of
Fibrovascular Ingrowth into Hydroxyapatite and Porous Polyethylene Orbital
Implants,’’ Ophthal. Plast. Reconstr. Surg.
, 10 [2] 96 (1994).
34
L. L. Hench, R. J. Splinter, W. C. Allen, and T. K. Greenlee, ‘‘Bonding
Mechanisms at the Interface of Ceramic Prosthetic Materials,’’ J. Biomed. Mater.
Res. Symp.
, 2 [1] 117–41 (1971).
35
L. L. Hench and J. Wilson, ‘‘Surface Active Biomaterials,’’ Science , 226 , 630–6
(1984).
36
P. de Wolff Technisch Physische Dienst, Delft, the Netherlands, ICCD
Grants-in-Aid (1957).
37
H. A. Batal, M. A. Azooz, E. M. A. Khalil, A. Soltan Monem, and Y. M.
Hamdy, ‘‘Characterization of Some Bioglass–Ceramics,’’ Mater. Chem. Phys.
, 80
[3] 599–609 (2003).
38
A. El Ghannam, E. Hamazawy, and A. Yehia, ‘‘Effect of Thermal Treatment on Bioactive Glass Microstructure, Corrosion Behavior, Potential, and Protein
Adsorption,’’ J. Biomed. Mater. Res.
, 55 [3] 387–98 (2001).
39
X. Chatzistavrou, T. Zorba, E. Kontonasaki, K. Chrissafis, P. Koidis, and K.
M. Paraskevopoulos, ‘‘Following Bioactive Glass Behavior Beyond Melting Temperature by Thermal and Optical Methods,’’ Phys. Status Solidi (a) , 201 [5] 944–
51 (2004).

S94 Journal of the American Ceramic Society—Mallick
40
D. C. Clupper, J. J. Mecholsky Jr., G. P. LaTorre, and D. C. Greenspan,
‘‘Sintering Temperature Effects on the In Vitro Bioactive Response of Tape Cast and Sintered Bioactive Glass–Ceramic in Tris Buffer,’’ J. Biomed. Mater. Res.
, 57
[4] 532–40 (2001).
41
D. C. Clupper, J. J. Mecholsky Jr., G. P. LaTorre, and D. C. Greenspan,
‘‘Bioactivity of Tape Cast and Sintered Bioactive Glass–Ceramic in Simulated
Body Fluid,’’ Biomaterials , 23 [12] 2599–606 (2002).
42
O. P. Filho, G. P. LaTorre, and L. L. Hench, ‘‘Effect of Crystallization on
Apatite-Layer Formation of Bioactive Glass 45S5,’’ J. Biomed. Mater. Res.
, 30 ,
509 (1996).
43
Q. Z. Chen, I. R. Thomson, and A. R. Boccacinni, ‘‘45S5 Bioglass s
-Derived
Glass–Ceramic Scaffold for Bone Tissue Engineering,’’ Biomaterials , 27 [11] 2414–
25 (2006).
44
I. K. Jun, Y. H. Koh, and H. E. Kim, ‘‘Fabrication of Highly Porous Bioactive Glass–Ceramic Scaffold with High Surface Area and Strength,’’ J. Am.
Ceram. Soc.
, 89 [1] 391–4 (2006).
45
P. Li, Q. Yang, F. Zhang, and T. Kokubo, ‘‘The Effect of Residual Glassy
Phase in a Bioactive Glass–Ceramic on the Formation of its Surface Apatite Layer
In Vitro,’’ J. Mater. Sci : Mater. Med.
, 3 [6] 452–6 (1992).
46
O. Peitl, E. D. Zanotto, and L. L. Hench, ‘‘Highly Bioactive P
2
O
5
–
Na
2
O–CaO–SiO
2
Glass Ceramics,’’ J. Non-Cryst. Solids , 292 [1–3] 115–26
(2001).
47
D. C. Clupper and L. L. Hench, ‘‘Crystallization Kinetics of Tape Cast Bioactive Glass 45S5,’’ J. Non-Cryst. Solid , 318 [1–2] 43–8 (2003).
48
A. E. Clark and L. L. Hench, ‘‘Calcium Phosphate Formation on Sol–Gel
Derived Bioactive Glasses,’’ J. Biomed. Res.
, 28 , 693–8 (1994).
49
L. L. Hench, ‘‘Sol–Gel Materials for Bioceramic Applications,’’ Curr. Opin.
Solid State Mater. Sci.
, 2 , 604–10 (1997).
50
W. Cao and L. L. Hench, ‘‘Bioactive Materials,’’ Ceram. Int.
, 22 , 493–507
(1996).
51
52
T. Kokubo, ‘‘Novel Bioactive Materials,’’ An. Quıˆm.
, 93 , S49–55 (1997).
V. Banchet, E. Jallot, J. Michel, L. Wortham, D. Laurent-Maquin, and G.
Balossier, ‘‘X-ray Microanalysis in STEM of Short-Term Physicochemical Reactions at Bioactive Glass Particle/Biological Fluid Interface.
Determination of O/Si Atomic Ratios,’’ Surf. Interface Anal.
, 36 [7] 658–65
(2004).
Vol. 92, No. S1
53
T. Kokubo, H. Kushitani, and S. Sakka, ‘‘Solutions Able to Reproduce In
Vivo Surface-Structure Changes in Bioactive Glass–Ceramic A-W3,’’ J. Biomed.
Mater. Res.
, 24 [6] 721–34 (1990).
54
C. Wang, T. Kashihiro, and M. Nogami, ‘‘Macroporous Calcium Phosphate
Glass–Ceramic Prepared by Two-Step Pressing Technique and Using Sucrose as a
Pore Former,’’ J. Mater. Sci.: Mater. Med.
, 16 [8] 739–44 (2005).
55
M. Kochs, C. H. Ko¨rber, B. Nunner, and I. Heschel, ‘‘The Influence of the
Freezing Process on Vapour Transport during Sublimation in Vacuum Freeze
Drying,’’ Int. J. Heat Mass Transfer , 34 [9] 2395–408 (1991).
56
K. A. Keler, G. M. Mehrotra, and R. J. Kernas, ‘‘Freeze Forming of Alumina
Monoliths’’; pp. 557–67 in Processing and Fabrication of Advanced Materials V ,
Edited by T. S. Srivatsan, and J. J. Moore. Minerals, Metals and Materials Society/AIME, Warrendale, PA, 1996.
57
S. W. Sofie and F. Dogan, ‘‘Freeze Casting of Aqueous Alumina Slurries with
Glycerol,’’ J. Am. Ceram. Soc.
, 84 [7] 1459–64 (2001).
58
Q. Fu, M. N. Rahaman, F. Dogan, and B. S. Bal, ‘‘Freeze Casting of Porous
Hydroxyapatite Scaffolds: I. Processing and General Microstructure Solutions,’’ J.
Biomed. Mater. Res. Part B: Appl. Biomater.
, 86B , 125–35 (2008), DOI: 10.1002/ jbm.b30997.
59
Q. Fu, M. N. Rahaman, F. Dogan, and B. S. Bal, ‘‘Freeze Casting of Hydroxyapatite Scaffolds for Bone Tissue Engineering Applications,’’ Biomed. Mater.
, 3
[2] 25005 (2008).
60
B.-H. Yoon, Y.-H. Koh, C.-S. Park, and H.-E. Kim, ‘‘Generation of Large
Pore Channels for Bone Tissue Engineering Using Camphene-Based Freeze Casting,’’ J. Am. Ceram. Soc.
, 90 [6] 1744–52 (2007).
61
E. R. Rubenstein and M. E. Glicksman, ‘‘Dendritic Growth Kinetics and
Structure II. Camphene,’’ J. Cryst. Growth , 112 , 97–110 (1991).
62
K. Anselme, ‘‘Osteoblast Adhesion in Biomaterials,’’ Biomaterials , 21 [7] 667–
81 (2000).
63
K. Lu, C. S. Kessler, and R. M. Davis, ‘‘Optimization of a Nanoparticle Suspension for Freeze Casting,’’ J. Am. Ceram. Soc.
, 89 [8] 2459–65 (2006).
64
S. Deville, E. Saiz, and A. P. Tomsia, ‘‘Ice-Templated Porous Alumina Structures,’’ Acta Mater.
, 55 , 1965–74 (2007).
65
T. Moritz and H.-J. Richter, ‘‘Ceramic Bodies with Complex Geometries and
Ceramic Shells by Freeze Casting Using Ice as a Mold Material,’’ J. Am. Ceram.
Soc.
, 89 [8] 2394–98 (2006).
&