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Journal of the European Ceramic Society 34 (2014) 2443–2452

Preparation of high open porosity ceramic foams via direct foaming moldedand dried at room temperature

Li-yuan Zhang, Da-li Zhou ∗, Ying Chen, Bin Liang, Jia-bei ZhouCollege of Materials Science and Engineering, Sichuan University, Chengdu 610064, China

Received 29 October 2013; received in revised form 17 January 2014; accepted 2 February 2014Available online 28 February 2014

bstract

erein an alternative approach was considered for addressing one difficulty of ceramic foams that the foam slurry with a high content of bubbleshich were obtained via direct foaming, cannot maintain well for a long time at room temperature. It is fascinating that the foam slurry mentioned

bove could stably mold and dry at room temperature, based on an animal protein as foaming agent, kaolin, talc powder and alumina as rawaterials, alpha-tricalcium phosphate prepared via co-precipitation as curing agent, and hydrophobic activated carbon powders as stabilizing

gent. Effects of the calcination temperatures, the contents of alpha-tricalcium phosphate and activated carbon powder on microstructures, crystalhases, compressive strength and open porosities of ceramic foams were studied systematically. The results indicated that ceramic foams with

high open porosity and uniform pore distribution and sizes sought for application in catalysts supports, could be produced by adjusting thesearameters.

2014 Elsevier Ltd. All rights reserved.

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eywords: Direct foaming; Alpha-tricalcium phosphate; Hydrophobic activate

. Introduction

Ceramic foams, one of the most important materials, haveeen widely used in a range of fields, namely filters, catalystsupports and bone scaffolds, due to their characteristics of highermeability, high porosity and specific surface area.1–3 In thesextensive applications, microstructural features of ceramicoams, for instance, the morphology, porosity, and pore size,re crucial factors that dominate the functional properties of thenal product. Ceramic foams were advantageous in terms ofreat property of low density,4 controllable microstructure,5–8

idely adjustable strength,9,10 high temperature resistance11,12

nd chemically inert composition. Direct foaming which is onef the most versatile among the methods to fabricate ceramicoams, offers a simple, low-cost and effective way to produce

eramic foams by incorporation of air bubbles into a suspensionf powders, and the pore structure of ceramic foams can beffectively controlled. However, the foam slurry cannot exist

∗ Corresponding author. Tel.: +86 28 85410272; fax: +86 28 85416050.E-mail address: zdl@scu.edu.cn (D.-l. Zhou).

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ttp://dx.doi.org/10.1016/j.jeurceramsoc.2014.02.001955-2219/© 2014 Elsevier Ltd. All rights reserved.

on powder; Room temperature drying

teadily for a long time at room temperature, derived fromhe thermodynamically unstable bubbles. Especially the foamlurry system with a high content of bubbles is prone to collapsend crack at room temperature, which will seriously affect theerformance of the final product. To date, freeze drying13,14

nd gelfoaming15 have been proposed to deal with the problem.owever, freeze drying is not suitable for mass production, alsoith a shortcoming of high cost. Gelfoaming which uses a largeumber of organic monomers and crosslinking agents, suffersrom high cost and environmentally hazardous behavior owingo the burning out of organics. In this work, a novel method wasonsidered for addressing this difficulty. Ceramic foams with aigh open porosity and uniform pore distribution and sizes, wererepared via direct foaming, based on an animal protein (cattleoof shell protein) as foaming agent, kaolin, talc powders andlumina as raw materials, alpha-tricalcium phosphate (�-TCP)repared via co-precipitation as curing agent, and hydrophobicctivated carbon powder as stabilizing agent.

Among animal protein foaming agents, egg white protein

EWP)16 and whey protein isolate (WPI)17 are the most fre-uent. Problems arising in Europe in the 1990s, due to theresence of animals in the feed chain of beef cattle, have raised

2444 L.-y. Zhang et al. / Journal of the European C

Table 1The mass ratios of chemicals to synthesize cattle hoof shell foaming agent.

Chemicals Mass ratios

Cattle hoof shell powder 63Sodium bisulfite 2.5Distilled water 630Calcium oxide 6.7Acetic acid Several dropsST

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odium dodecyl benzene sulfonate 7.8riethanolamine 3.2

ocio-economic concerns and questions regarding the use ofhe blood collected in slaughterhouses.18 Remadnia et al.18 hassed atomized bovine hemoglobin derived from animal bloods foaming agent to manufacture cementitious concrete com-osites. However, it is rarely reported that cattle hoof shell wassed as a foaming agent to prepare ceramic foams. Concerning

foam stabilizing agent, particles are the most popular.8,19–21

ivaldini et al.19 has analyzed why foams containing col-oidal hydrophilic particles are unstable, and Mishra et al.21

as prepared porous silica tiles, involving the use of partiallyydrophobized silica powder particles for stabilization of foams.o the best of our knowledge, hydrophobic activated carbonowder was firstly used in the present work as a foam stabilizinggent.

. Experimental

.1. Synthesis of cattle hoof shell foaming agent

The cattle hoof shells were immersed and boiled in dilutelkali solution for 4 h, and then dried in an oven for 24 h at0 ◦C. The previously dried samples were crushed, followedy ball-milling for 12 h with a rotational speed of 180 rpm and

medium of zirconium ball to aim powder particles with ca.0 �m in diameter. To produce the liquid of protein foaminggent, NaHSO3 and cattle hoof shell powders were added toistilled water at 80 ◦C for 4 h. The temperature was controlledy a water bath. Then the reaction mixture was heat-treated for

h at 100 ◦C after the addition of CaO. After that, the mix-ure was cooled and filtered, following acetic acid was addedo adjust pH to 7. Finally, sodium dodecyl benzene sulfonate

nd triethanolamine were added to the solution, stirred gently.able 1 summarizes the mass ratios of chemicals to synthesizeattle hoof shell foaming agent.

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able 2he weight/volume details of chemicals to fabricate ceramic foams.

hemicals Solid weights (wt.%)

-TCP 3, 5, 10, 15, 20ctivated carbon powders 3, 5, 10, 15, 20aolin, talc powder and alumina 100 − W�-TCP − WC

olyvinyl alcohol (PVA) solution 36.35

ixture of Na2HPO4, NaH2PO4 and citric acid solution 36.35

attle hoof shell foaming agent 9.1istilled water 18.2

s/Vl = 15/11, where, Ws, Vl is the total weight of the solids (g), the total volume of

eramic Society 34 (2014) 2443–2452

.2. Preparation of alpha-tricalcium phosphate (α-TCP)

Alpha-tricalcium phosphate (�-TCP) was prepared via co-recipitation by controlling Ca/P in molar ratio 3/2. A solutionf Ca(NO3)24H2O was slowly added to the mixed solution ofH4HCO3 and (NH4)2HPO4 at a water bath 37 ◦C, followedy using concentrated ammonia water to adjust pH to 7. Pre-oma powders were obtained after vacuum filtration, washinghe precipitate exhaustively with anhydrous ethanol, and dry-ng. The presoma powders were firstly calcined at 1260 ◦C for

h and then taken out from the furnace immediately, follow-ng simultaneously agitated to cool down to room temperature.ventually, fine particles of �-TCP were produced by planetaryall milling for 3 h with a rotational speed of 180 rpm and mediaf anhydrous ethanol and zirconium ball.

.3. Fabrication of ceramic foams

The mixture of kaolin, talc powders and alumina, �-TCPnd activated carbon powders were mixed uniformly in the pro-ortions given in Table 2. Then a certain amount of polyvinyllcohol (PVA, 1788 ± 50) solution and a mixed solution whichonsisted of Na2HPO4, NaH2PO4 and citric acid, were added toain ceramic slurry, following stirred evenly. Foams were pre-ared with cattle hoof shell foaming agent and distilled wateria direct foaming. Ceramic foam slurry system was obtained bydding ceramic slurry to foams, agitated homogeneously. Sub-equently, ceramic foam slurry was poured into a cylindricalold which was coated a fine layer of petrolatum previously, to

orm a cylinder-shaped, followed by drying at room temperatureo gain ceramic foam green body. The samples were demoldeduring the drying process. The calcinations of ceramic foamreen bodies were carried out at 1050–1150 ◦C with 25 ◦C innterval for 2 h in air, followed by cooling down to room tem-erature in furnace. The weight/volume details of chemicals toabricate ceramic foams are shown in Table 2. During the sus-ension preparation, some points deserved to be mentioned areiven below.

Alpha-tricalcium phosphate was chosen for curing agentased on our previous work.22 A mixed solution of Na2HPO4,aH2PO4 and citric acid was added to promote the hydration

8,19–21

eaction of �-TCP. Consulting the reports that particlesave been used as foam stabilizing agents, the hydrophobic acti-ated carbon powder was used as a foam stabilizing agent owingo its advantages of hydrophobicity, low density, porous surface

/Liquid volumes (vol.%) Ratios/concentrations

In mass ratios 19/8/39.1 wt.%0.2 M, 0.2 M and 6.7 wt.%, respectively

the liquids (ml), respectively.

ean Ceramic Society 34 (2014) 2443–2452 2445

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CPaaipvtncf33r1eCto cristobalite phase23 owing to the formation of a large amountof dissociative SiO2 (Cristobalite, SiO2, JCPDS, CARD, No.39-1425). Diffraction peaks of anorthite occurred at 2θ of 23.8, 27.8,

L.-y. Zhang et al. / Journal of the Europ

tructure, and many strong adsorption sites. The concentrationf PVA solution we opted was 9.1 wt.%, because a higher vis-osity of the suspension could increase the surface strength ofhe foam films and retard the growth of the bubbles, resulting in

ore stable foams. A selection of kaolin/talc powder/aluminan mass ratios 19/8/3 was related to the following. It was foundhat a certain content of kaolin in raw materials gave an appropri-te strength to ceramic foams. However, a high enough contentf kaolin made an easy densification of the inner wall of theore, this being not benefit for adhesion of the catalysts. Aroper content of talcum powder could increase surface rough-ess of the ceramic foams, which was helpful for adsorption ofhe catalysts. Nevertheless, a high number of talcum powdersrought a low mechanical strength and an easy pulverization ofhe material. After a comprehensive consideration, kaolin, tal-um powder and alumina were adopted in mass ratios 19/8/3.ptimization of the mixture preparation was studied by chang-

ng the �-TCP and activated carbon powder contents, as well ashe calcination temperature, and keeping other parameters con-tant. When we evaluated the effect that the parameters brought,e put emphasis on observing no crack and collapse of the foam

lurry system, uniform pore distribution and sizes, as well as aigh open porosity.

.4. Characterization of materials

The crystalline phases of the samples were investigated by-ray diffract meter (DX-1000, Dangdong Fangyuan Instru-ent Co., Ltd., China) using Cu K� radiation with a scanning

ate of 0.06◦ s−1 between 10 and 70◦ and a working volt-ge/current of 40 kV/25 mA. The microstructures of ceramicoams were recorded by field emission scanning electronicroscopy (FESEM, S-4800, Hitachi, Japan). Samples forESEM analysis were sprayed some gold particles to enhance

he conductivity. Ceramic foams were cut into cylindricalieces (� 20 mm × 20 mm) and polished by alumina emeryaper prior to compressive strength test. The average com-ressive strength based on five samples were determined usingn universal test machine (XDL-50000, Jiangdu Vinporearesting machine factory, China) with a cross-head speed of0 mm min−1. Open porosities of ceramic foams were deter-ined by the Archimedes’ method using water as buoyantedium. The weight of the sample was measured by an elec-

ronic balance (AUY120, Shimadzu Corporation, Japan) with aensitivity of 10−4 g.

. Results and discussion

.1. Effects of α-TCP

Fig. 1 presents XRD pattern of �-TCP prepared via co-recipitation. The diffraction peaks of calcium phosphateatched very well with the standard card of �-TCP (Calcium

hosphate, Ca3(PO4)2, JCPDS, CARD, No.09-0348) indicat-

ng that calcium phosphate prepared via co-precipitation inhe present work was main �-TCP. However, diffraction peaksscribed to �-TCP (Calcium Phosphate, Ca3(PO4)2, JCPDS,

ig. 1. XRD pattern of �-TCP prepared via co-precipitation fired at 1260 C for h.

ARD, No.09-0169) and calcium pyrophosphate (Calciumyrophosphate, Ca2P2O7, JCPDS, CARD, No.09-0345) werelso detected. Fig. 2 represents SEM imagine of �-TCP. It waspparent that the �-TCP particle sizes were homogeneous, typ-cally 2–3 �m. Fig. 3 shows XRD patterns of ceramic foamsrepared from different contents of �-TCP and 5 wt.% acti-ated carbon powders treated at 1100 ◦C/2 h. Obviously, whenhe content of �-TCP was 3 and 5 wt.%, the major compo-ents of the samples were found to be very similar, mainlyonsisted of cordierite, cristobalite, mullite, as well as calciumeldspar. The characteristic peaks at 2θ of 16.4, 26.4, 31.0, 33.2,5.3, 39.3 and 40.9◦ corresponded to mullite phase (Mullite,Al2O3·SiO2, JCPDS, CARD, No.15-0776), resulted from theeaction between Al2O3 and SiO2. Diffraction peaks at 2θ of0.5, 18.2, 19.0, 21.7, 26.4, 28.5, 29.5, 33.4 and 38.5◦ weressentially due to cordierite (Cordierite, Mg2Al4Si5O18, JCPDS,ARD, No. 13-0294). One strong peak at about 21.8◦ indexed

Fig. 2. SEM imagine of �-TCP.

2446 L.-y. Zhang et al. / Journal of the European C

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ig. 3. XRD patterns of ceramic foams prepared from different contents of-TCP and 5 wt.% activated carbon powders calcined at 1100 ◦C for 2 h.

5.7 and 43.1◦ (Anorthite, CaAl2Si2O8, JCPDS, CARD, No.41-486), as a result of the interactions among �-TCP, mullite andristobalite. At the content of 10 wt.%, the diffraction inten-ity of cristobalite at 21.8◦ was lowered, while that of calciumeldspar was increased, in the absence of mullite, this behav-or being related to the formation of anorthite from mullite andristobalite. At the same time, increased in the diffraction inten-ities of cordierite was associated with the reactions among talc,ullite, cristobalite and alumina.Further improving the content to 15 wt.%, the crystal phases

howed significant changes with respect to 10 wt.%. The diffrac-ion intensities of cordierite degraded sharply while that ofalcium feldspar increased. It was because the cordierite pro-ided Al, Si and O elements which were essential to formnorthite. As the initial content of �-TCP reached up to 20 wt.%,iffraction peaks of cordierite were almost disappeared whilehat of calcium feldspar increased further. In addition, it wasotable that the diffraction peak of cristobalite at 21.8◦ wasower than that of 15 wt.% �-TCP. On one hand, diffractioneak at about 21.8◦ was ascribed to cristobalite and cordieritet 15 wt.% �-TCP. Transformation from cordierite into anor-hite occurred when the content of �-TCP increased from 15 to0 wt.%. On the other hand, cristobalite was required to supplyi element to form anorthite due to the fact that cordierite couldot provide enough Si when the content of cordierite was low.

Fig. 4 gives the SEM images of ceramic foams obtained bydding 5 wt.% activated carbon powders and �-TCP contentsarying from 3, 5, 10, 15 to 20 wt.% treated at 1100 ◦C/2 h.s shown, when the initial content of �-TCP was only 3 wt.%,

he structures in the longitudinal section of the samples wereot uniform. Distribution of the top pores was dense with aper-ures ∼100 �m while that of the bottom ones was extraordinarilyneven. A broad distribution of the bottom pore sizes was also

bserved. The non-uniformity of the top and bottom pores dis-ributions was improved at the content of 5 wt.%. With rising theontent of �-TCP up to 10 wt.%, distribution of the pores in the

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eramic Society 34 (2014) 2443–2452

ongitudinal section was uniform with pore sizes in 100–200 �m.owever, variations of the homogeneous pore distribution and

izes were observed when the content further reached 20 wt.%.ome large pores were more than 500 �m in diameter. Theeason responded for that is given below.

Bubbles obtained from a protein foaming agent via directoaming were in a thermodynamically unstable state as a con-equence of foam decay. Foams processed by direct foamingf slurries have a tendency to develop defects such as cracks,elamination, warpage during drying.24 Usually, the foam slurryystem with a high content of bubbles is prone to collapse andrack during drying at room temperature. It is generally acceptedhat there are two main mechanisms of foam decay, i.e. therainage of liquid films and the gas diffusion permeating liq-id films. There is a direct relationship among the two kindsf mechanisms and the liquid film properties, as well as thenteraction between liquid film and Plateau boundary.25 Therainage of liquid films is a result of the mutual extrusion ofubbles and the influence of gravity. The influence of gravity onrainage depends on the capability of interface slipping,26 whilehe gas diffusion permeating liquid films mainly comes from theon-uniformity of the bubble sizes.

After molding of the foam slurry with a high content ofubbles, in the absence of �-TCP and activated carbon pow-ers, the pressure the bottom bubbles received was significantlyreater than the top ones because of the gravity force resultingn easier rupture of the bottom bubbles. Accordingly, the foamlurry flowed, and the whole system collapsed accompanied byhe bursting of the bubbles. However, by adding an appropriateontent of �-TCP, hydroxyapatite (HA) acicular crystal coulde formed through the hydration reaction of �-TCP.22 A retic-lar structure (Fig. 4h) could be further observed through therosslink of the acicular crystals.

Due to repelling by the reticular structure, on one hand, thenterface slipping became more difficult so that effect of theravity force on the drainage of the liquid films was reduced;n the other hand, the flowing speed of the foam slurry wasecelerated. In addition, the growth of the bubbles was retarded.esides, this function was helpful to weaken inhomogeneousffect of the gravity force on the bubbles. The reticular structureshat �-TCP provided through hydration reaction was few whenhe content of �-TCP was quite low. Hence, rupture of the bottomubbles occurred owing to effect of the gravity. The rupture timef the bubbles was different, followed by unevenly flowing ofhe foam slurry under the performance of the gravity.

As a result, asymmetrical distribution and sizes of the bot-om pores were found. Whereas, with the extension of time, aertain pressure could inhibit the growth of the bubbles, whichas responsible for homogeneous distribution and apertures of

he top pores. However, water was the main component of theoaming agent and hence the foam films. When the content of-TCP was exceeded a critical value, the hydration reaction of-TCP brought a strong absorption of moisture from liquid films

neven pore distribution and sizes were caused.Fig. 5 illustrates the variations of compressive strength and

pen porosities of ceramic foams as a function of �-TCP (5 wt.%

L.-y. Zhang et al. / Journal of the European Ceramic Society 34 (2014) 2443–2452 2447

Fig. 4. SEM images of ceramic foams and green body prepared from 5 wt.% activated carbon powders and different contents of �-TCP. (a) 3 wt.%, 1100 ◦C; (b)5 wt.%, 1100 ◦C; (c) 10 wt.%, 1100 ◦C, low magnification; (d) 10 wt.%, 1100 ◦C, high magnification; (e) 15 wt.%, 1100 ◦C, low magnification; (f) 15 wt.%, 1100 ◦C,h ◦ ◦ pore.

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igh magnification; (g) 20 wt.%, 1100 C; (h) 10 wt.%, 25 C, inner wall of the

ctivated carbon powders) calcined at 1100 ◦C/2 h. As indicated, negative correlation between compressive strength and openorosities was revealed in ceramic foams. With the increasinghe content of �-TCP, the compressive strength of the samples

educed initially, increased then, and degraded finally, while thepen porosities performed contrarily, this being probably relatedo the following.

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As the content of �-TCP increased from 3 to 5 wt.%, the majoromponents of the mixture were found to be similar, however,he pore sizes increased obviously. While carrying out mechan-cal test on the samples with bigger pore sizes, the area which

ould be effectively contacted with plunger tip on the surfaceot smaller. Whereby, a lower stress would be used to damagehe samples under the condition of loading force.27 Besides, this

2448 L.-y. Zhang et al. / Journal of the European Ceramic Society 34 (2014) 2443–2452

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ig. 5. The variations of compressive strength and open porosities of ceramicoams as a function of �-TCP (5 wt.% activated carbon powders) treated at100 ◦C for 2 h.

henomenon could be ascribed to the poor integrity of the porest 5 wt.% as well. The two reasons above both eventually gave anncreased open porosity. Further rising the content to 10 wt.%,he pore structures were integrated, and the pore distributionas homogeneous, as well as pore sizes. As a consequence,

ncreased in the area which could be effectively contacted withlunger tip on the surface was brought. The loading force alsoould be evenly dispersed in the specimens. All that resulted inn increase in compressive strength. Similarly, the pores with rel-tively integrate structure were isolated from each other, whichas connected with the decrease of the open porosity. With the

ontent of �-TCP further improving from 10 to 15 and 20 wt.%,he excessive �-TCP absorbed an amount of moisture fromhe liquid films due to curing behavior, leading to fracture of

any foams. Hence, the integrity of the pores was reduced, andhe pore size distribution was extremely asymmetrical. Besides,ubbles united each other resulting in an increment of the per-eability (open porosity) and a decrement of the corresponding

ompressive strength, although the inner wall of the pore in thepecimens that added 15 wt.% �-TCP was more compact thanhat of 10 wt.% (Fig. 4d and f).

.2. Effect of activated carbon powder

Our previous work showed that surface of the foam slurryystem with a high content of bubbles was still easy to cracky only adding �-TCP without activated carbon powders, whichould seriously affect the performance of ceramic products. The

apillary stress rise due to evaporation of water from the surfacelays a role in formation of cracks.28,29 When the foam slurryith a high content of bubbles was molded and dried in contactith air, the surface area was easy to dry first, even to form lumps.o, the capillary became smaller in diameter, or even blocked.onsequently, the internal moisture was difficult to evaporate

hrough the capillary increasing the humidity difference betweenhe internal and external. Hence, the surface layer received the

ensile stress, while the inner layer born the compressive stress,ollowed by easy cracking of the surface. In addition, duringhe drying process, the subsidence of the solids originated from

tfit

Fig. 6. A SEM image of the surface of the activated carbon powder.

ffect of the gravity force, decreased the content of solid phasen the surface layer. Due to this fact, the distance between theolid particles was apparently increased, also followed by easyracking of the surface.

In this work, consulting the reports8,19–21 that particles haveeen used as foam stabilizing agents, the hydrophobic activatedarbon powder was used as a foam stabilizing agent to dealith the problem owing to its advantages of hydrophobicity, lowensity, porous surface structure, and many strong adsorptionites (Fig. 6). As expected, good results have been observed.

Figs. 7a, 4c, 7d and e reveal the SEM images of ceramic foamsbtained by 10 wt.% �-TCP and contents of activated carbonowder varying from 3, 5, 10 to 20 wt.% treated at 1100 ◦C/2 h.otably, pores were disintegrated, and distribution of the pore

izes was broad when the content of activated carbon powderas only 3 wt.%. At the content of 5 wt.%, distribution and

izes of the pores in the longitudinal section were uniform, withpertures in 100–200 �m and strut thicknesses about 2–10 �mFig. 7c). With further improving the content to 10 wt.%, distri-ution and sizes of the pores were homogenous with obviouslyncreased apertures, while the integrity of the pore structure waseduced. As the content of activated carbon powder achieved0 wt.%, the integrity of the pores reduced apparently, and manytruts cracked resulting many pores coalesced (Fig. 7e). Thehenomenon can be elucidated by the reasons given below.

The hydrophobic activated carbon powders would suspend inhe suspension owing to their low density. Some carbon powdersurrounded the bubbles acting against foam coalescence, whilehe others were adsorbed to a gas–liquid interface.

It is known that when a particle is attached to a gas–liquidnterface, the overall Gibbs free energy of the system is reduced,esulting in a more favorable thermodynamic condition. How-ver, the adsorption energy of a particle at gas–liquid interfacebubble’s interface) must be at least few hundred times greaterhan the thermal energy in order to allow foam stabilization.19

therwise the thermal energy could withdraw particles from thenterface of the bubbles.19 Due to the hydrophobicity, a high con-

act angle between the activated carbon powders and the foamlms was formed resulting in great adsorption energy. Hence,

he powders did not have enough energy to break loose from

L.-y. Zhang et al. / Journal of the European Ceramic Society 34 (2014) 2443–2452 2449

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ig. 7. SEM images of ceramic foams prepared from 10 wt.% �-TCP and diffagnification; (b) 3 wt.%, high magnification; (c) 5 wt.%; (d) 10 wt.%; (e) 20 w

he gas–liquid interface, which was associated with more stableoams.

Moreover, the activated carbon powders had porous surfacetructures and many strong adsorption sites, making the lamellarggregate particles adsorbed on the surface by a mode of physicaldsorption. With a synergistic effect of �-TCP and activated car-on powders, it was more difficult for the solids and the bubbleso move in the suspension. Decreased in the capillary diame-er also was harder, improving channels and surface areas fromhich the inner water could move outward. Besides, the decre-ent of the solid loading in the upper foam slurry was also found

o be more difficult. Accordingly, easy cracking of the surfaceayer was restricted, and uniform structures were obtained.

However, as reported by Prud’homme et al.,30 gas evolutioned to foamed macrostructures. Activated carbon powders wereurned at a temperature ca. 464 ◦C, at which the strength of theeramic body was low. When the content of activated carbon

owder was high enough, the strength of the pore walls wasot enough to counterbalance the gas pressure6 owing to theurning out of a large number of activated carbon powders in

oto

contents of activated carbon powder fired at 1100 C for 2 h. (a) 3 wt.%, lowow magnification; (f) 20 wt.%, high magnification.

he calcining process. Thus anomalistic pore morphology, crack,nd even the collapsed ceramic body were observed (Fig. 7e and).

The contribution of activated carbon powders for the com-ressive strength and open porosities of ceramic foams waslso taken into account (10 wt.% �-TCP), as given in Fig. 8.s shown, the compressive strength and open porosities sub-

tantially exhibited a negative relation. With the increase ofctivated carbon powders, the compressive strength of theamples increased firstly and then decreased, while the openorosities performed contrarily. When the contents of �-TCPnd activated carbon powder both were 10 wt.%, the compres-ive strength and open porosity of ceramic foams treated at100 ◦C/2 h were 1.37 MPa and 78.8%, respectively. With theontent of activated carbon powder increasing from 3 to 5 wt.%,he pore distribution and sizes were more even so that the loadingorce could be well distributed in the samples under the condition

f loading force.27 So, a higher force would be used to destroyhe specimens. It was worth noting that the compressive strengthf the samples was quite high (4.05 MPa) when the content was

2450 L.-y. Zhang et al. / Journal of the European Ceramic Society 34 (2014) 2443–2452

Fig. 8. Compressive strength and open porosities of ceramic foams preparedfa

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Fig. 11 plots the compressive strength and open porosities ofsamples treated at temperatures ranging from 1050 to 1150 ◦C,taken every 25 ◦C (10 wt.% �-TCP and 5 wt.% activated carbon

rom 10 wt.% �-TCP and different contents of activated carbon powder treatedt 1100 ◦C for 2 h.

nly 3 wt.% owing to the formation of acicular structure whichorresponded to mullite phase,31 as shown in Fig. 7b. Furthermproving the content to 10 wt.%, the integrities of the poretructures decreased, and the connectivity between the pores,s well as pore sizes, increased notably. Hence, increased in thepen porosities and decreased in the corresponding compressivetrength were given.

It should be noted that the increase of open porosity was notpparent when the content of carbon improved from 10 to 15 andurther to 20 wt.%. One possible explanation to this phenomenonas that the small pores that the activated carbon powders left in

he original position closed or disappeared after the thermal treat-ent, this being probably related to densification of the innerall of the pore. The compressive strength overall experienced

reduction due to the propagation of cracks in the specimensFig. 7f).

.3. Effect of calcination temperature

The crystalline phases as a function of temperature are exhib-ted in Fig. 9 (10 wt.% �-TCP and 5 wt.% activated carbonowders). At 1050 ◦C some reflections of cristobalite, anor-hite were detected, as well as mullite. Characteristic peak atθ of 21.8◦ was associated with cristobalite phase. Two weakeaks at 23.8, 24.5◦ and two diffraction peaks at 27.8, 35.7◦ndexed to anorthite. Additionally, no peaks of cordierite but

ullite were detected. With the temperature rising up to 1075 ◦C,he components changed significantly. The diffraction intensi-ies of anorthite peaks enhanced while that of mullite decreased.esides, the characteristic peaks of cordierite were also detected.

t was worth mentioning that Rodrigues Neto and Moreno23

lso used kaolin, talc and alumina as raw materials, and theyid not detect the signal of cordierite until the reaction tem-erature increased up to 1300 ◦C (for the unmilled sample).

bove-mentioned discussion indicated that the addition of �-CP and activated carbon powders could effectively degrade

he formation temperature of cordierite phase.

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ig. 9. XRD patterns of ceramic foams calcined at different temperatures10 wt.% �-TCP and 5 wt.% activated carbon powders).

In order to further figure out what produced reduction in theormation temperature of the cordierite, �-TCP, activated carbonowders, or �-TCP and activated carbon powders? Two otherxperiments (one for addition of 10 wt.% �-TCP, the other for

wt.% activated carbon powders) were performed. The resultsFig. 10) showed that cordierite phase could also be detected byddition of 10 wt.% �-TCP, while no reflections of cordieritehase were detected by 5 wt.% activated carbon powders. Itould be demonstrated that it was the addition of �-TCP caus-ng a lower formation temperature of cordierite phase rather thanctivated carbon powders.

ig. 10. XRD patterns of ceramic foams obtained by adding 10 wt.% �-TCP, wt.% activated carbon powders, and 10 wt.% �-TCP and 5 wt.% activatedarbon powders treated at 1075 ◦C for 2 h.

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influence of surfactant concentration. Mater Chem Phys 2009;113:441–4.7. Medri V, Ruffini A. The influence of process parameters on in situ inorganic

foaming of alkali-bonded SiC based foams. Ceram Int 2012;38:3351–9.

ig. 11. Compressive strength and open porosities of ceramic foams calcined atifferent temperatures (10 wt.% �-TCP and 5 wt.% activated carbon powders).

owders). In most cases, the strength of the ceramic prod-cts gets higher along with the increasing of the calcinationemperature. Here a contrast tendency occurred in this work.bviously, with the increase the sintering temperature, the com-ressive strength of the ceramic foams increased initially andhen degraded, while the open porosity decreased firstly andmproved late. However, the minimum open porosity and theighest compressive strength did not experience at the same tem-erature, indicating that the increment of compressive strengthas not only depended on the decrease of open porosity. With

he temperature rising up from 1050, 1075 to 1100 ◦C, theompressive strength progressively increased from 3.24, 4.38o 4.41 MPa, respectively. The main causes responded for thatere, on one hand, with the improving the calcination tempera-

ure, the relative content of cordierite, anorthite and cristobaliten the specimens increased, which gave a higher mechanicaltrength; on the other hand, pore sizes became smaller because ofhe shrinkage. Besides, this phenomenon could also be attributedo homogeneous distribution and sizes of the pores in ceramicoams prepared from 10 wt.% �-TCP and 5 wt.% activated car-on powders calcined at 1100 ◦C/2 h (Fig. 4c). However, theompressive strength greatly reduced from 3.20 to 1.93 MPaith the temperature further increasing from 1125 to 1150 ◦C.ne reason was, the shrinkage of pores increased, as the cal-

ination temperature grew. During this process, the shrinkageercentages of pores might not be completely consistent, fol-owing the pore distribution and sizes were more and moresymmetrical, and the shape of pores became more and morenomalistic that from the approximate sphericity to irregularolygon. The loading force could not be evenly dispersed in theamples, leading to some cracks from stress concentration. Thether was, pore struts became thinner, even fractured with theurther rise of the temperature, which eventually contributed to

deterioration of the compressive strength.As the temperature rose up from 1050 to 1125 ◦C, open

orosities reduced continuously from 75.3%, 73.8%, 71.8% to

2.7%, respectively. Nevertheless, the open porosity increasedp to 70.2% when the calcination temperature further reached150 ◦C. It was because in a certain temperature range, with

eramic Society 34 (2014) 2443–2452 2451

he improving of temperature, the specimens shrunk to a certainxtent which did not lead to fracture of the pore struts resultingn initial decrease of the open porosity. However, fracture of theore struts occurred when the temperature further reached upo 1150 ◦C, leading to the improved connectivity between theores and hence, the open porosity increased.

. Conclusions

Based on a protein (cattle hoof shell protein) as foaminggent, kaolin, talc powder and alumina as raw materials via directoaming, foam slurry with a high content of bubbles could moldnd dry well at room temperature by adding �-tricalcium phos-hate (�-TCP) prepared via co-precipitation as curing agent,nd hydrophobic activated carbon powder as stabilizing agent.eramic foams with a high open porosity and homogeneousore distribution and sizes, could be produced by optimizing thealcination temperature, the contents of �-TCP and activatedarbon powder. The parameters mentioned above affected theicrostructures, crystal phases, compressive strength and open

orosities of ceramic foams significantly. The calcination tem-erature in formation of cordierite had decreased greatly afterhe addition of �-TCP (1075 ◦C in the present work). Whenhe contents of �-TCP and activated carbon powder both were0 wt.%, the foam slurry with a high content of bubbles couldtably mold and dry at room temperature, and the compres-ive strength and open porosity of ceramic foams treated at100 ◦C/2 h were 1.37 MPa and 78.8%, respectively.

cknowledgements

Thanks are due to Jin-chun Shi and Yong Jin for the usefulharacterizations of the samples, and to Hai-tao Zhang and Liou for the revising of the paper.

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