Znd. Eng. Chem. Res. 1994,33, 860-870
860
Improvement of Thermal Stability of Porous Nanostructured Ceramic Membranes Y ue-Sheng Lin,' Chih-Hung Chang, and Ramakrishnan Gopalan Department of Chemical Engineering, University of Cincinnati, Cincinnati, Ohio 45221
Unsupported crack-free lanthana-doped alumina and titania membranes and yttria-doped zirconia membrane were prepared by the sol-gel method. A novel solution-sol doping method was employed to coat the dopant oxide on the grain surface of these nanostructured ceramic membranes. The pore and phase structure data of these membranes after heat treatment from 450 to 1100 "C under an atmosphere of air and a steam/air mixture show a substantially improved thermal and hydrothermal stability of these doped ceramic membranes. Doping lanthana (in y-alumina and titania membranes) and yttria (in zirconia membrane) raises the y-alumina to a-alumina, anatase to rutile, and tetragonal zirconia to monoclinic zirconia phase transformation temperature by about 200 "C (for alumina), 150 "C (for titania) and 300 "C (for zirconia), respectively. At temperatures lower than the phase transformation temperatures, doping retards the surface area loss and pore growth of the three membranes. For the three ceramic membranes investigated, the effects of stabilizing the pore structure decrease in the following order: zirconia > titania > y-alumina.
Introduction Nanostructured alumina, titania, and zirconia membranes prepared by the sol-gel method are finding applications in liquid separation and filtration (Burggraaf et al., 1989). They also offer potential applications in high temperature catalytic reactions and gas separations (Egan, 1989; Bhave et al., 1991). In these high temperature applications,the membranesare usedas supports on which a gas separation layer (e.g., dense oxide or metallic thin layer, microporous silica layer) or catalyst layer is coated. These nanostructured ceramic membranes might also directly be employed in hot gas filtration applications (Alvin et al., 1991). Thus, thermal and hydrothermal stability of these membranes are important to their potential high temperature applications. The thermal stability of aceramic membrane is referred to the ability of the membrane to withstand a prolonged heat treatment under a particular atmosphere. A membrane being thermally stable at a certain temperature means that the properties (phase structure, mechanical strength, and, most importantly, the pore structure) of the membrane remain unchanged or changed negligibly a t that temperature for a period of time comparableto the practical application time. Among these three nanostructured ceramic membranes, the y-A1203 membrane is by far the best studied one. Burggraaf and his co-workers in the early 1980s investigated systematically the synthesis of unsupported and supported y-Al203 membrane by the sol-gel method. The details on the characteristics of boehmite particles, membrane micropore structures, dip-coating process, and membrane properties were reported in a series of publications by Burggraaf and his co-workers (Leenaars et al., 1984;Leenaars and Burggraaf, 1985a,b,c). Subsequently, more studies on the sol-gel preparation and properties of the y-alumina membrane were reported by the groups of Burggraaf, Cot, and Anderson, and other groups worldwide (Larbot et al., 1988; Anderson et al., 1988; Okubo et al., 1990; Cini et al., 1991; Uhlhorn et al., 1989, 1992a,b). Understanding of the synthesis and gas permeation properties of the alumina membrane has been significantly
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improved as a result of these investigations. Titania and zirconia membranes with a pore size in the range of 3-5 nm were also prepared by several investigators using the sol-gel method (Anderson et al., 1988;Larbot et al., 1989a; Gieselman et al., 1988;Lijzenga et al., 1991). However, in comparison with data for aluminaand titania membranes, information on the sol-gel preparation of the zirconia membrane is rather limited (Chang et al., 1993). High temperature properties of these porous ceramic membranes have attracted considerable interest because of the potential high temperature applications of these membranes. Burggraaf and his colleagues (Leenaars and Burggraaf, 1985a; Burggraaf et al., 1989) have earlier determined the pore size of some ceramic membrane toplayers after heat-treatmenta (in air) at different temperatures. They reported that the pore size of the y-alumina top-layer increases significantly at a temperature range of 900-1000 "C. A same set of data of the pore size vs sintering temperature for y-alumina, titania, and zirconia membranes (in air, 1 h sintering time) was reported by Cot and his co-workers in several publications (Larbot et al., 1988,1989a,b). In a study on the calcination and the thermal stability of alumina membrane, van Veen et al. (1989) investigated the pore structure of the y-alumina membrane as a function of firing temperature in a relatively low temperature range (425-600 "C). The increase in the pore diameter of the studied membranes was found to be less than 1 nm after 800 h of heat treatment in the temperature range mentioned. Lin and his co-workers (Chang et al., 1993) recently performed a comprehensive comparative study on the thermal and hydrothermal stability of all the three nanostructured (alumina, zirconia, and titania) membranes. They found that for a fixed firing time (30 h), phase transformation of these three ceramic membranes occurs in the following temperature ranges: 900-1100 "C for an alumina membrane (y-Al203 to (Y-A~~OS), 600-900 "C for a zirconia membrane (tetragonal to monoclinic) and 450-700 "C for a titaniamembrane (anatase to rutile). In these temperature ranges, phase transformation results in a significant change in the membrane pore structure of these three membranes due to the formation of larger crystals of the new phase. Pore sizes of the alumina, zirconia and titania membranes increase roughly by factors of 15, 3 and 2, respectively, after the completion of the
OSSS-5SS5/94/2633-0S60~04.50/00 1994 American Chemical Society
Ind. Eng. Chem. Res., Vol. 33, No. 4,1994 861 phase transformation. Furthermore, the new phase has a different lattice structure and volume from those of the original phase. For a supported membrane, the phase transformation can cause cracks because of the change in the lattice volume. The pore structure of these three ceramic membranes changes, to a lesser extent, after firing at temperatures below the phase transformation temperature ranges. Sintering by surface diffusion is primarily responsible for the change of the pore structure (Chang et al., 1993).The sintering results in a decrease in surface area and an increase in the pore size of these ceramic membranes. Among these three ceramic membranes investigated, the rate of the pore structure changes as a result of sintering decreases in the following order: zirconia > titania > alumina. Consideringsintering and phase transformation, the alumina membrane is the most thermally stable one, while the titania is the least stable. Lin and co-workers (Chang et al., 1993)also found that presence of steam in the atmosphere, which is often encountered in industrial applications, accelerates the rate of the pore structure change of these three ceramic membranes at elevated temperatures. The surface diffusion (sintering) and nucleation and crystal growth (phase transformation) rates are enhanced by the presence of steam in the heat-treatment atmosphere. The increase of the pore size of the membranes after firing under the steam/ air atmosphere is about 50-300% more than that of the same membranes fired at the same temperature under the atmosphere of air. Although ceramic membranes are known to be much more thermally stable than polymeric membranes, the results summarized above show that the pore structures of these three membranes change a t elevated temperature as a result of sintering and phase transformation of alumina, titania, and zirconia. The practical applications of these ceramic membranes at high temperatures depend on, among others, the thermal and hydrothermal stability of these membranes. In this work, a grain surface coating technique originally developed by Lin et al. (1991)was extended to improve the thermal and hydrothermal stability of alumina, zirconia, and titania membranes prepared by the sol-gel method. Such an improvement could result in a new group of ceramic membranes with a longer lifetime and a higher application temperature. The grain surface coating technique will have application not only in improving the thermal stability of ceramic membranes but also in modifying the surface chemistry of other advanced materials fabricated by the sol-gel method. The main objective of this paper is to report the results of the comparative study on the thermal stability improvement of these three ceramic membranes in unsupported form by the grain surface coating technique. The preparation and properties of these ceramic membranes in supported form will be reported in a separate paper. It should be pointed out that the study on the unsupported membranes is important because (1)unsupported ceramic membranes are easier to characterize, allowinga quicker development of a new method for property improvement, and (2) preparation of crack-free unsupported membranes would help identify proper composition and concentration of sols from which crack-free supported membranes can be successfullyfabricated (Lin et al., 1989;Lin and Burggraaf, 1991).
Strategy for Thermal Stability Improvement Sintering and transformation from a metastable phase to stable phase of the alumina, zirconia, and titania
membranes at elevated temperatures are thermodynamic spontaneous processes (Chang et al., 1993). It was found that sintering of crystalline y-alumina, zirconia (tetragonal phase), and titania (anatase) membranes is dominated by the surface diffusion mechanism (Chang et al., 1993).In this case, the sintering rate is proportional to the specific surface energy of the crystallites (Shi et al., 1991;Chang et al., 1993). Hence, covering the surface of a metal oxide grain (first oxide) by an oxide with a lower sintering rate (secondoxide) will lower the surface energy of the particles, thus reducing the sintering rate of the material. Chang et al. (1993)also found that phase transformation of the three ceramic membranes proceeds via the mechanism of nucleation (on grain surface or boundary) and crystal growth. The phase transformation rate is proportional to the number of nucleation sites on the grain surface (boundary) and activation energy for nucleation (Hishita, 1983;Chang et al., 1993). Retardation of the phase transformation of a metal oxide (the first oxide) by doping a second oxide is possible if the dopant covers the grain surface of the first oxide, thus reducing the grain boundary nucleation sites and/or increasing the activation energy for nucleation (Hishita, 1983; Schaper, 1984; Mercera, 1991). The improvement of the thermal stability of these membranes can be achieved with a strategy aimed at reducing the specific surface energy and the number of nucleation sites, or increasing the activation energy for sintering and phase transformation. Monolayer coating of a second oxide (as catalyst) on the surface of the first oxide particles (as support) has been an active topic of research in the catalysis field (e.g., Lipsch and Schuit, 1969;Schaper, 1984;Mercera, 1991;Bettman et al., 1989; Tijburg et al., 1991). I t is found that the formation of the monolayer of the second oxide on the surface of the first oxide particles is a thermodynamically favorable process if the second oxide, with a metal ion size larger than the size of the metal ion of the first oxide, can form a strong surface bond with the first oxide particles (Xie and Tang, 1990). The formation of the monolayer of MOOS,CuC12, VZO~, La~03,etc. (ascatalysts) on the surface of y-alumina or silica particles (as support) has been confirmed experimentally (Xie and Tang, 1990). In this work, the thermal and hydrothermal stability of alumina, titania, and zirconia membranes was improved by doping lanthanum oxide or yttrium oxide on the grain surface of these membranes prepared by the sol-gel method. Lanthana (for alumina and titania) and yttria (for zirconia) were chosen as the second oxide for the following reasons: (1)the size of La3+ is larger than Al3+ and Ti4+,and the radius of Y3+ is 1.06 A, also larger than the Zr4+ radius (0.87A) (Shi et al., 1991);(2) it has been shown that lanthana and yttria are effective in stabilizing structure of alumina and zirconia catalyst supports (Schaper, 1984;Mercera, 1991;Shi et al., 1991). The other reason for selecting yttria as the second oxide for coating the zirconia membrane was due to the consideration of the potential application of a yttria-doped zirconia membrane as a base support for solid oxide fuel cells and dense oxygen semipermeable ceramic membranes. In the catalysis field, coating an oxide monolayer on the surface of catalyst supports is normally conducted by the solid-dispersion method and wet-impregnation method (Xie and Tang, 1990). It is however difficult to use these methods to dope a controlled amount of the second oxide in the supported top-layers of these ceramic membranes prepared by the sol-gel method. In this work, a solutionsol mixing method reported earlier by Lin et al. (1991)
862 Ind. Eng. Chem. Res., Vol. 33, No. 4, 1994
was employed to dope the second oxide into the ceramic membranes. As will be shown later, this novel doping method allows coating of a two-dimensional surface layer (possibly a monolayer) of the second oxide on the grain surface of the three ceramic membranes.
Experimental Section Preparation of the lanthana- and yttria-coated porous ceramic membranes included the following steps: (1) synthesizing the stable boehmite, titania and zirconia sols; (2) preparing La(N03)3and Y(NOd3 solutions with a controlled pH value; (3) mixing the lanthanum nitrate solution with each of the three oxide sols in a proper composition; (4) pouring a controlled amount of the lanthanum- or yttrium-doped sol into a petri dish and converting the sol to xerogel by controlled drying; (5) calcining the unsupported gels to form ceramic membranes. The 1M boehmite, 0.25 M titania, and 0.25 M zirconia sols were prepared by hydrolysis and condensation of aluminum tri-sec-butoxide (Janssen), titanium tetraisopropoxide (Aldrich) and zirconium n-propoxide (Alfa), respectively. PVA (polyvinyl alcohol) solution and HPC (hydroxypropyl cellulose) solution were prepared and used as the drying control chemical additive (DCCA) for preventing crack formation during the drying process. Details on the preparation of these three oxide sols and PVA and HPC solutions are given elsewhere (Chang et al., 1993). The 0.3 M lanthanum nitrate solution was synthesized by dissolving 5.4 g of La203 (Fluka) into 100 mL of 1 M HN03 solution at 50 "C. The solution was filtered using filtering paper to obtain pure salt solution. The 0.07 M yttrium nitrate solution was prepared by mixing 2 g of Y(N03)3 (Alfa) with 95 mL of water and 5 mL of 1 M HN03at 50 "C. The pH values of the yttrium nitrate and lanthanum nitrate solutions were 1 and 3 respectively. Alumina, titania, and zirconia membranes which were not doped with the second oxide (referred to as undoped membranes) were prepared from the boehmite sol (mixed with PVA), titania sol (mixed with PVA and HPC), and zirconia sol (mixed with PVA). The starting sols for synthesizing the lanthana- or yttria-doped membranes (referred to as doped membranes) were the same boehmite, titania, and zirconia sols (with PVA and/or HPC) used for synthesizing the undoped membranes. Lanthana and yttria were doped in the alumina, titania, and zirconia membranes using the following solution-sol mixing method. A given amount of the lanthanum nitrate or yttrium nitrate solution was thoroughly mixed with a controlled amount of each of the three oxide sols (doped with DCCA). Mixing lanthanum or yttrium nitrate directly with the sol allows an accurate control of the final molar ratio of lanthana or yttria to alumina, titania, or zirconia. Volumes of the sols, the DCCA solution, and the nitrate solution used to prepare the unsupported membrane samples are summarized in Table 1. The atomic ratio of the metal of the dopant oxide to the metal of the membrane oxide is also given in the table. In doping lanthanum or yttrium nitrate in the oxide sols, the pH value of the mixed sol was kept essentially unchanged so that the doped sols were still stable with a negligible change in particle size. The isoelectric points of the beohmite, titania, and zirconia sols are around pH of 5-6. In this work, the pH values of the boehmite, zirconia, and titania sols were respectively 2.5, 1.5, and 0.5. The pH values of the lanthanum or yttrium nitrate
Table 1. Volumes (mL) of Oxide Sols, DCCA and Nitrate Solutions Used in Preparing Unsupported Membrane Samples membrane oxide sol PVA HPC nitratea atomic ratiob 0 2 0.03 T-Alz03 20 13 38 1 0.07 Ti02 16 20 ZrOz 14 6 0 1 0.02 ~~~
a La(N03)3 for alumina and titania membranes; Y(NOd3 for the zirconia membrane. b La/A1 for the alumina membrane, La/Ti for the titania membrane, Y/Zr for the zirconia membrane.
doped sols were essentially same as compared with those of the corresponding undoped sols. Unsupported alumina, titania, or zirconia membranes of about 20 pm thick and about 80 mm in diameter (both doped and undoped with lanthanum or yttrium) were prepared by drying a given amount of sol in a drying oven at 40 OC and 40-5096 relative humidity. The dried xerogels were calcined in a temperature programmable furnace at 450 OC for 3 h, with carefully controlled heating and cooling rates (Chang, 1993). It should be noticed that PVA and/ or HPC were used in both the undoped and lanthana- or yttria-doped membranes. These experimental procedures were based on the following understanding of the sol-gel processing. The original undoped particulate sols contained suspended aggregatesconsistingof many primary particles. Solutionsol mixing introduced the salts into the surrounding of the aggregates. The salt ions were adsorbed on the surface of the primary particles during the aging and/or drying step. The metal nitrate molecules were converted to the corresponding metal oxide, which remained on the surface of primary particles, after calcination at 450 OC under an atmosphere of air for 3 h. Typical undoped alumina, zirconia, and titania membranes were obtained from the sols prepared respectively by mixing 20 mL of beohmite sol with 13 mL of PVA solution (30 g/L) (for alumina membrane), 14 mL of zirconia sol with 6 mL of PVA (30 g/L) (for zirconia membrane), and 16.4 mL of titania sol with 20 mL of PVA solution (1g/L) and 38 mL of HPC solution (3.5 g/L) (for titania membrane). The lanthana- or yttria-doped alumina, zirconia, and titania membranes were prepared by mixing the PVA- and/or HPC-doped sols specified above with about 1-2 mL of the lanthanum or yttrium nitrate solution. To investigate the thermal and hydrothermal stability improvement of these ceramic membranes, samples of the undoped and doped membranes were heat-treated under the same conditions. Heat treatments were done by placing the samples in the temperature programmable furnace at different temperatures (500-1100 "C) under an atmosphere of (1)air and (2) a steam/air mixture (molar ratio 1:l). Most samples were fired for 30 h. Firing experiments in the atmosphere of steam/air were performed with an atmospheric retort which is a gastight, heavy gauge steel enclosure placed inside the regular box furnace. The details on this heat treatment equipment under humid atmosphere are reported elsewhere (Chang et al., 1993). The surface area, pore volume, and pore size distribution (average pore size) of the unsupported membranes were characterized by Nz ad(de)sorption isotherms using adsorption porosimeter (Micromeritics, ASAP 2000). The phase structures of the ceramic membranes were identified by an X-ray diffractometer (Siemens, Cu Kal). The phase compositions and the crystallite sizes of the zirconia and titania membranes were calculated from the XRD reflections of these membrane samples. The weight
Ind. Eng. Chem. Res., Vol. 33, No. 4, 1994 863 fraction of rutile present in the titania membranes was calculated using the equation (Spurr and Myers, 1957)
w,= 1/[1+ 0.81~(101)/1~(110)]
(1)
where IA(101) and I ~ ( 1 1 0 )are the X-ray integrated intensities of the (101) reflection of anatase and the (110) reflection of rutile, respectively. The weight fraction was converted to the volume fraction using the density of the titania phases. The volume fraction ( V,) of the monoclinic zirconia phase was calculated by (Toraya et al., 1984)
V , = 1.311Xm/[1+ O.311Xm1
2.4
1
AIz03 _____ LaU n dDo op pe de d AI2O3
1
1.6
1.6
7
I
(2)
with the integrated intensity ratio X,
x,
= [z,(iii)
+ ~ , ( i i i ) i / [ ~ , ( i i i )+ z,(iii) + 1,(iii)i
(3) where Z,(111) and Z,(llT) are t_he X-ray integrated intensities of the (111) and (111) reflections of the monoclinic zirconia phase and It(lll)is the intensity of the (111)reflection of the tetragonal phase. The crystallite sizes of the tetragonal and monoclinic zirconia, anatase, and rutile titania were calculated with the Scherrer equation (Cullity, 1978) using the (101) reflection of anatase titania, the (110) of rutile titania, (111) of tetragonal zirconia and (111) and (111) of monoclinic zirconia (Spurr and Myers, 1957; Toraya et al., 1984).
Results and Discussion Preparation and Properties of Doped Membranes. Membrane Formation. Stable undoped and nitratedoped beohmite, zirconia and titania sols were obtained by the experimental procedures described above. After the drying and calcination steps, these sols were converted to crack-free unsupported ceramic membranes. Serious cracks were observed on the lanthana- or yttria-doped ceramic membranes, especially yttria-doped zirconia membrane, when they were not prepared under the optimum conditions specified above. The pore size distributions of lanthanum-doped alumina and titania and yttrium-doped zirconia membranes are compared in Figure 1 with those of the corresponding undoped membranes. The pore structure data of these six membranes are summarized in Table 2. For the alumina, titania, and zirconia membranes, doping of the second oxide results in a slight reduction in the pore volume. After doping of the second oxide, the average pore size increases and pore size distribution broadens for the alumina membrane, whereas the opposite changes are found for the titania and zirconia membranes. It is known that the pore structure (measured by the nitrogen adsorption method) is determined by the size and shape of the primary particles and drying and calcination conditions (which are identical for the doped and undoped membranes). The slight difference in the pore structures between the undoped and oxide-doped ceramic membranes is more likely due to the change in the size and shape of the primary particles in the sols after being doped with the salt solution. The introduction of the salt ions (La3+ or Y3+) and the change in Nosconcentrations in boehmite, titania, and zirconia solsmight alter slightly the charge distribution around the sol particles, affecting the shape and size of the primary particles. Nevertheless, the structural differences between the undoped and doped membranes are small from the viewpoint of practical application. These comparison results show that the doping method employed in this
10
1
0.32
I
I
-Y
______
1
D o p e d Zr02 U n d o p e d ZrO2
10
Pore Diameter ( n m )
Figure 1. Pore sizedistributions of lanthana- or yttria-doped yAl203, TiO2, and ZrO2 membranes (after calcination at 450 OC for 3 h).
study has negligible effects not only on the sol stability but also on the pore structure of the membranes. It was reported that among these three oxide sols, stability of zirconia sol and formation of crack-free zirconia gel were most sensitive to the sol preparation conditions (Chang et al., 1993). Thus, extreme care was taken in doping Y(N03)3 and the organic binder in stable zirconia sol as slight change in the pH, ion concentration, and other conditions could easily destablize the sol. It was found that the stability of the zirconia sol could be maintained after doping Y (NO313with a ratio of yttrium to zirconium not larger than 5 % . The optimum PVA/ZrOz-sol ratio was found to be 30 % as far as the sol stability and formation of crack-free gel are concerned. Precipitation of large zirconia particles in the sol or formation of severe cracks in the zirconia xerogel were observed when the zirconia sols were not prepared under these optimum conditions. The XRD patterns of the lanthana- or yttria-doped alumina, titania, and zirconia membranes are compared with the correspondingundoped ones in Figure 2. ?-A1203 (cubic phase) typically has very weak XRD reflections (Wefers and Mirsa, 1987). As shown in Figure 2a, no sharp reflection peaks are observed for both the lanthana-doped and undoped alumina membranes. Nevertheless,the XRD patterns for the lanthana-doped and undoped ?-A1203 membrane are identical. The XRD patterns for the undoped zirconia and titania membranes indicate a tetragonal phase structure for both materials. Essentially identical XRD patterns (including 28 values) are found for the yttria-doped and undoped zirconia membranes, as shown in Table 3 and Figure 2c. The XRD patterns of the lanthana-doped titania membrane after calcination does not exhibit clearly the reflection peaks of the anatase phase, as shown in Figure 2b. It appears that the doped titania membrane comprises
864 Ind. Eng. Chem. Res., Vol. 33, No. 4,1994 Table 2. Comparison of Pore Structures of Doped and Undoped y-A12O)a,TiOz, and ZrO2 Membranes (Calcined at 460 OC for 3 h) ?-&Os Ti02 ZrO2 undoped La-doped undoped La-doped undoped Y-doped d, nm 3.0 3.1 6.8 5.8 3.5 3.1 339 290 138 159 66.7 66.1 SBET, m2/g V,, mL/g 0.30 0.28 0.37 0.32 0.13 0.11 I
I
a
I
I
A40, Membranes ( 4 ° C ) Doped TiO,
7
b
I
I
io, Membmnes (4SOT-3
700°C
600°C
450°C
28 (deg) Figure 3. XRD patterns of lanthana-doped titaniamembranes after heat treatments at different temperatures.
I
0
J'O
40
50
6'0
20 (deg) Figure 2. Comparison of XRD patterns of lanthana-doped alumina (a) and titania (b) and yttria-doped zirconia (c) membranes with those of undoped membranes (after calcination at 450 OC for 3 h). The phase structures are identified as anatase for the undoped titania membrane (b)and tetragonal for both the doped and undoped zirconia membranes (c). Table 3. Comparison of 28 Values (deg) of Major XRD Reflection Peaks between Yttria-Doped and Undoped Zirconia Membranes 20 value reflection of tetragonal-Zr02 undoped Y-doped (111) 30.37 30.36 (202) 50.43 50.49 (220) 50.64 50.67 (131) 60.25 59.94 (200) 35.22 35.14
acertain amount of amorphous titania. The XRDpatterns of the doped titania membrane after heat treatment at higher temperatures are given in Figure 3. The anatase reflections,although observable,are very broad. The XRD pattern for the doped titania membrane after heat
treatment at 900 "C already shows clearly the rutile reflections. The titania xerogel after drying was in amorphous form. For the undoped titania membrane the calcination step (450"C for 3 h) not only consolidated the titania gel but also facilitated the phase transformation from amorphous titania to tetragonal titania (anatase). For the doped titania membrane, the XRD results given in Figures 2b and 3 indicated that the heat treatment at different temperatures does not completely convert the amorphous titania to anatase. This seems to suggest doping lanthana retards the phase transformation from amorphous titania to tetragonal titania. The XRD data show that no X-ray detectible phases other than those of y-alumina, titania (anatase), or zirconia (tetragonal) are present in the lanthana- or yttria-doped membranes. This indicates that no crystals of the dopant (La203or Y203) with XRD detectible size are formed in the membranes. These results suggest that dopant oxide is either present in a form of two-dimensional layer (most likely a monolayer) on the grain surface of these three ceramic membranes (Xie and Tang, 1990)or incorporates into the lattice structure of the alumina, zirconia, and titania. Identification of the exact state of the dopant oxide in the nanoscale particles is rather difficult. Since the two-dimensional layer model has been often used in catalysis literature (e.g., Lipsch and Schuit, 1969;Schaper, 1984;Mercera, 1991;Bettman et al., 1989;Xie and Tang, 1990;Tijburg et al., 1991)to explain the results of doping oxide on catalyst support, this model is also adopted here to explain the experimental findings reported next. Pore and Phase Structure Evolution. Figure 4shows the surface areas of the three lanthana- or yttria-doped
Ind. Eng. Chem. Res., Vol. 33, No. 4, 1994 865 350
Doshed curves -- steam/oir Solid curves -- oir
--
--
0'4 0.3 h
3 E!
i Dashed curves Solid curves
steam/air
air
-
m
1 '. E
200
v
s
0.2
-
0
$1 2
0.1
4
300
700 900 Firing t e m perat u r e ("C)
500
1100
Firing temperature ("C)
Figure4. Effectaoffiiingtemperatureandatmosphereonthesurface area of the lanthana- or yttria-doped ceramic membranes (30-hheat treatment time).
-- stearn/oir -- air
Doshed curves Solid curves
- 0 300
500 700 900 Firing temperature ("C)
o-30,o
I
1100
Figure 5. Effects of firing temperature and atmosphere on the average pore size of the lanthana- or yttria-doped ceramicmembranes (30-h heat treatment time).
alumina, titania, and zirconia membranes after heat treatments at various temperatures for 30 h under air or steam/air atmosphere. In all cases, the surface area decreases with increasing firing temperature. The larger the initial membrane surface area, the more the reduction in the surface area for that membrane. At the same temperature, the presence of steam in the firing atmosphere enhances the surface area reduction. For these three membranes, this enhancing effect on surface area reduction increases with the increasing initial surface area of the membrane. Figures 5 and 6 show the average pore size and pore volume of the three lanthana- or yttriadoped ceramic membranes after heat treatments at the same conditions as for the results reported in Figure 4. For all three lanthana- or yttria-doped membranes, the average pore size increases with increasing firing temperature. At temperatures lower than 900 "C, the firing temperature has the strongest effect on the change of the average pore size of the titania membrane and the least effect on the alumina membrane. For most samples, the pore volume increases and, after reaching a maximum, decreases with increasing firing temperature. In all cases, the presence of steam enhances the.pore size increase, as shown in Figure 5. However, the effects of steam on the pore volume is somewhat complicated. For alumina, the presence of steam slows the pore volume shrinking rate, as shown in Figure 6 where the dashed curve (through open triangle symbols) is higher than the corresponding solid curve at temperature lower than 900
Figure 6. Effects of firing temperature and atmosphere on the pore volume of the lanthana- or yttria-doped ceramic membranes (30-h heat treatment time).
"C. For both the titania and zirconia membranes, the presence of steam enhances the pore volume shrinking rate. As will be shown later, observable phase transformations for the doped y-alumina (y to a), zirconia (tetragonal to monoclinic), and titania (anatase to rutile) start around 1100,900, and 600 O C , respectively (referred to the phase transformation starting temperature). It should be noted that y-alumina actually transforms to a-alumina via 6-alumina (around 800 "C) and &alumina (around 900 "C) (Wefers and Misra, 1989). Similar to y-alumina, both &alumina and &alumina have very weak XRD reflections and are difficult to identify by XRD. In the discussion, 6-alumina and &alumina are included in y-alumina. Within the heat treatment temperature ranges investigated, both sintering process (below the phase transformation starting temperature) and phase transformation (above the phase transformation starting temperature) are responsible for the surface area reduction and pore structure change. Chang et al. (1993) found that sintering of the undoped alumina, zirconia, and titania membranes is dominated by the surface diffusion mechanism, of which the surface area reduction can be described by (German, 1977)
ASIS = k [ y t e ~ p ( - E / R l ' ) / T l ~ / ~ . ~ (4) where S is the surface area, y is the specific surface energy, E is the activation energy for surface diffusion, T and t are the temperature and time of heat treatment, and k is a constant. Thisequation shows indeed that the reduction in the surface area, AS, is roughly proportional to the initial surface area. The enhancing effecb of steam on surface area reduction is a result of increasing specific surface energy or reducing activation energy for sintering. For membranes of crystalline ceramic materials, surface area reduction accompanies the pore size increase due to the removal of materials from the convex (or flat) grain surface to the more concave necks. This phenomenon was explained in detail for the sintering of the undoped alumina, zirconia and titania membranes (Chang et al., 1993). It is known that sintering should result in pore volume reduction. Results given in Figure 6, however, show that in most cases the pore volume increases and, after reaching a maximum, decreases with increasing firing temperature. The seeminglyabnormal results are believed to be associated with the use of large molecular drying control chemical additives (DCCA) (PVA and HPC). These molecules, after burn off, caused microstructure nonuni-
866 Ind. Eng. Chem. Res., Vol. 33, No. 4,1994
formity which might be responsible for the initial increase of the pore volume (Lin et al., 1991;Chang et al., 1993). For pure alumina, zirconia,and titania membranes without doping PVA and/or HPC, the pore volume decreases monotonically with increasing firing temperature (Chang et al., 1993). For the nano-structured ceramic membranes, phase transformation proceed via nucleation on the grain surface (boundary) and crystal growth (Chang et al., 1993). Thus, the formation of large grains of the new phase accompanies a further reduction in surface area and increase in pore size. The effects of the sintering and phase transformation on the change of the pore structure of these lanthana- or yttria-doped ceramic membranes can in general be explained by the surface diffusion mechanism (for sintering) and surface nucleation and crystal growth mechanism (for phase transformation). The overall effects of the heat treatment on the pore structure change are similar between the undoped and the lanthana- or yttria-doped ceramic membranes. Nevertheless, the comparison presented next will show a significant improvement of the pore structure stability and many other interesting features of these lanthana- or yttria-doped ceramic membranes. Comparison w i t h Undoped Membranes. Phase Transformation. XRD patterns of the lanthana- or yttria-doped alumina, zirconia, and titania membranes after heat treatments at 1100 "C (alumina) or 900 "C (zirconia and titania) for 30 h are compared in Figure 7 with those of the corresponding undoped ceramic membranes. As shown in Figure 7a, the undoped alumina membrane after this heat treatment has already converted to 100% a-alumina phase, as compared with the initial y-alumina structure (see Figure 2a). No observable a-A1203 phase is found in the lanthana-doped alumina membrane after the same heat treatment. Doping lanthana has at least raised the phase transformation starting temperature for y-A1203by 200 "C (from 900to 1100 "C). Similarly, parts b and c of Figure 7 show that undoped titania and zirconia membranes after these heat treatments have transformed from the initial phases (anatase and tetragonal zirconia, see Figure 2b,c) to their thermodynamically more stable phase (rutile and monoclinic zirconia). The lanthana-doped titania membrane after heat treatment at 900 "C for 30 h consists primarily of the rutile phase (see Figure 7b). But the yttria-doped zirconia membrane remains essentially in the tetragonal phase after the same heat treatment. The retardation on the phase transformation due to doping with the second oxide is better illustrated by the volume fraction data shown in Figure 8. These data were calculated from the XRD reflections of the membranes after heat treatments at different temperatures. Figure 8 shows that considerable amount of rutile is found in the undoped titania membrane at 450 "C. Roughly about 20 % monoclinic zirconia is formed in the undoped zirconia membrane at 600 "C. In comparison, phase transformation of the doped titania and zirconia membranes starts at much higher temperatures. For the doped membranes, the temperatures requires to achieve the same amount of conversion to the more stable phase (e.g., 20% rutile for titania and 10% monoclinic for zirconia) are respectively about 300 "C (for zirconia) and 150 "C (for titania) higher than that of the corresponding undoped membrane. In other words, doping yttria or lanthana has raised the titania and zirconia phase transformation temperature by about 150 and 300 "C, respectively. Phase transformation from y-alumina to a-alumina,
I
I
a
I
I
Membranes (llOO"0
I
1
I
I
I
TiO,Membranes (9oOV
& Undoped
20 ( d 4
Figure 7. Comparison of XRD patterns of the lanthana- or yttriadoped alumina, titania, and zirconia membranes with the undoped membranes after heat treatment under air atmosphere at a high temperature for 30 h. The undoped alumina, titania, and zirconia membranes are in the form of a-alumina, rutile, and monoclinic(with small amount of tetragonal), respectively.
anatase to rutile, and tetragonal to monoclinic zirconia is thermodynamically spontaneous a t the temperature range investigated (metastable to stable phase transformation) (Wefers and Misra, 1987;Wilson, 1979;Vahldiek, 1966; Scott, 1975; Garvie, 1978). Small grains of tetragonal zirconia could be stabilized a t room temperature when the grain size is smaller than 30 nm (Garvie, 1978). There are no thermodynamic equilibrium phase transformation temperatures for metastable to stable phase transformation processes, which are controlled by the kinetic factors. Lattice structure change requires sufficient activation energy, which is overcome by raising the sample temperature. On the other hand, formation of a new solid phase in nanoparticles usually proceeds via the nucleation and crystal growth process. Therefore, nucleation sites are needed to facilitate the phase transformation. The temperature at which new stable phase is formed from a metastable phase depends on the lattice structure difference between the two phase and nature of the nucleation sites and the properties of grain surface. In general, the rate of the formation of new phase (on volume basis) depends on the nucleation and crystal growth rates, which are determined by the nucleation sites and activation energy for nucleation and crystal growth (Christian, 1975).
Ind. Eng. Chem. Res., Vol. 33, No. 4, 1994 867 .-6
1 .o
I
7 0.8
'1
- O"
6
0.6
?
0.4
'1 0.0
CUU
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:pedq
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900
Doped Zr02
0.0
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,
Kn
d -Steam/Air - -- - ---
A A A M Undoped AIIO, +LLu La Doped A120J4
-A7-;*;Q 1 1
,
t
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6.0 I
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hU4b Undoped AIIO
,dAAAAA Undoped Ti02
,
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c
,
0.0
St:%?(*-
b
A-!
a
I
L
9
1 7
a
I I
4 1
, , ,
500 700 900 Firing temperature ("C)
1100
Figure 8. Comparison of the volume fraction of rutile in lanthanadoped titania membrane (a), and of the monoclinic phase in yttriadoped zirconiamembrane (b) withthat of the correspondingundoped membranes after heat-treatment under air atmosphere at different temperatures (30 h).
On the basis of the surface nucleation mechanism (the nucleation sites are the defects present on the grain surface or boundary), the volume fraction of the new phase, x , can be correlated to the heat-treatment temperature, T, and time, t, by the Avrami equation (Christian, 1975; Brinker and Scherer, 1990)
i 0
'
0 3 h 0 " : 0 -03 Firing t e m p e r a t u r e ("C)
Figure 9. Comparison of the effects of heat treatment on the pore structure of the lanthana-doped alumina membrane (solid curves, close symbols) and that of the undoped alumina membrane (dashed curves, open symbols) SOB^ and do are the surface area and average pore size given in Table 2). 1.o 1
x = 1- exp(-(kt)")
Firing t e m p e r a t u r e ("C)
(5)
birA M M U n d o p e d TO -La Doped Ti&
0.8
with
k = k, exp(-EIRT) (6) where E is the activation energy for nucleation, R is the gas constant, and k, and n are constants, with n = 1 for the surface nucleation mechanism (Christian, 1975). A combination of eqs 5 and 6 gives ln[ln(l - x ) ] = ln(k,t)
- E/RT
p
-0.6
0.6
?
A
-0.4
0.4 \
\
0.2
-0.2
A
1
(7)
0.0
If the effects of heating and cooling on the phase transformation are neglected, the activation energy for nucleation can be estimated from volume fraction data given in Figure 8 based on the Avrami plots. The activation energy estimated is 253 kJ/mol for the undoped titania membrane and 400 kJ/mol for the undoped zirconia membrane. These activation energy data are consistent with those measured by other methods (Chang et al., 1993). Because of insufficient volume fraction data in the temperature range investigated the activation energies for the doped titania and zirconia membranes were not estimated. Nevertheless, the higher phase transformation temperatures for the doped membranes as compared with the corresponding undoped ones indicate that doping with the second oxide raises the activation energy for both membranes. The increase in the phase transformation temperature and the reduction in the phase transformation rate is probably because the doped oxide covers the grain surface of these three membranes. The presence of a twodimensional layer on the grain surface may reduce the number of nucleation sites (defect or vacancy) and increase the activation energy for phase transformation, thus, lowering the overall phase transformation rate.
0.0
I
3 0
.
.
...
.
I
1
,
560 700 960 Firing temperature
.
I
0
.
0
1I b O 2
("C)
Firing temperature
("C)
Figure 10. Comparison of the effects of heat treatment on the pore structure of the lanthana-doped titaniamembrane (solidcurves,close symbols) and that of the undoped titania membrane (dashed curves, open symbols).
Pore Structure and Grain Size. Figures 9-11 compare the effects of firing temperature on the surface area and average pore size of the lanthana or yttria doped alumina, zirconia and titania membranes after heat treatments at different temperatures and atmospheres with those of the corresponding undoped membranes. As
868 Ind. Eng. Chem. Res., Vol. 33, No. 4, 1994 b
~1.0
s:jtL titania > alumina. The specific grain surface energy of these materials also decreases in this order. This is consistent with the surface diffusion model, as described by eq 4. For the same firing temperature, the surface area of the doped alumina, titania, and zirconia membranes is larger than that of the corresponding undoped membranes. The difference in the surface area between the doped and undoped membranes decreases in the following order: zirconia > titania > alumina. This is the same as the order of the sintering rate. As suggestedearlier, the second oxide doped is presented in the form of a two-dimensional layer (probably a monolayer) on the grain surface of the alumina, zirconia, and titania membranes. Formation of the two-dimensional
Ind. Eng. Chem. Res., Vol. 33, No. 4, 1994 869 required for the 90% tetragonal to monoclinic transformation at 900 and 600 "C. At 600 "C,the time required for a 20% tetragonal to monoclinic transformation is 30 h for the undoped zirconia membrane and about 1400 h for the doped zirconia membrane.
W Tetragonal Undoped oQWMonoclinic Undoped u*u Tetragonal Doped
Conclusions
" I
300
560
700
Firing temperature
900
1100
('C)
Figure 13. Crystallite size of tetragonal and monoclinic phases in zirconia membranes as a function of heat treatment temperature.
tion of the stable phase is not very clear, but it is obvious that the stable phase has a much larger grain size than that of the metastable phase. It is important to note that in all cases, both the metastable and stable phases of the doped membranes have a smaller grain size than the corresponding undoped membranes. For the alumina membrane the grain size can not be measured from XRD data, but it can be inferred from the pore structure variation data that considerable grain growth takes place starting from 700 "C. Finally, it should be mentioned that doping a second oxide cannot completely retard the structure change of these membranes at elevated temperatures. But the results presented above have clearly demonstrated that covering a second oxide on the grain surface of these nanostructured ceramic membranes reduces the sintering and phase transformation rates. The improvement on the stability of the pore structure is more significant if one considers it on the time scale. The change of the pore structure or composition due to sintering or phase transformation depends exponentially on the temperature but linearly on time, as shown by eqs 4 and 7 rewritten in the following forms: for sintering
[AS/S13.6 = Kt exp(-E/RT)
(8)
for phase transformation ln(1- x ) = -Kt exp(-EIRT)
(9)
K is a constant. For a certain extent of the change in the pore structure at a given period of time, the temperature required for the doped membrane may be only a few tens of degrees higher than that of the corresponding undoped one. However, a much longer time is required for the doped membrane as compared with the corresponding undoped membrane for the same extent of change in pore structure a t a lower temperature. This can be demonstrated using the following example. For the undoped and doped zirconia membranes, about 90% and 10% tetragonal zirconia has respectively transformed to monoclinic zirconia after the heat-treatment a t 900 "C for 30 h (see Figure 8b). With these data and the activation energy for phase transformation for the undoped zirconia membrane (570 kJ/mol), the activation energy for phase transformation for the doped zirconiamembrane, calculated using eq 9, is about 870 kJ/mol. I t can be estimated again using eq 9 that for the doped zirconia membrane roughly about 700 and 14000h are respectively
Unsupported lanthana-doped alumina and titania membranes and a yttria-doped zirconia membrane were prepared by a solution-sol mixing method. The pore structure of these lanthana-doped or yttria-doped membranes is similar to that of the corresponding undoped membranes. Optimum conditions were identified for preparing stable sols which gave crack-free unsupported lanthana- or yttria-doped membranes. These optimum conditions are important for fabrication of crack- and pinhole-free supported membranes doped with lanthana or yttria. The thermal and hydrothermal stability of these three ceramic membranes was significantly improved by doping lanthana or yttria. Doping lanthana (in y-alumina and titania membranes) and yttria (in zirconia membrane) raises the y- to a-alumina, anatase to rutile, and tetragonal to monoclinic phase transformation temperatures by 200 "C (for alumina), 150 "C (for titania) and 300 O C (for zirconia), respectively. Doping also retards the surface area loss and pore growth of the three membranes at temperatures lower than the phase transformation temperatures. The effects of stabilizing the pore structure decrease in the following order: zirconia > titania > y-alumina. Doped lanthana or yttria is possibly present in the form of a two-dimensional layer (or even monolayer) on the grain surface of these three ceramic membranes. On the basis of the mechanism of surface diffusion for sintering, the lanthana or yttria layer lowers the specificgrain surface energy of the y-alumina, titania, and zirconia membranes, thus reducing the sintering rate of the membranes. The mechanism of surface nucleation and crystal growth could explain the phase transformation process observed in these three ceramic membranes. Covering lanthana or yttria on the grain surface of the three membranes increases the activation energy for nucleation as well as reduces the number of nucleation sites, thus raising the phase transformation temperature and decreasing the phase transformation rate (on volume basis). This solution-sol mixing method for grain surface coating can in principle be applied for coating of materials on the grain surface of the particles prepared by the solgel method for the purpose of modifying the surface chemistry and other properties of the sol-gel-derived ceramic materials.
Acknowledgment This project was supported in part by the National Science Foundation (Grant CTS-9212272).
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