β-aluminas

Solid State Ionics 46 ( 1991 ) 319-324 North-Holland

Improved precursor ceramics for proton-conductive 13"/13-aluminas A i c h u n T a n , C h u K u n K u o a n d P a t r i c k S. N i c h o l s o n Ceramic Engineering Research Group, Department of Materials Science and Engineering, McMaster University, Hamilton, Ontario, Canada LSS 4L 7

Received 19 February 1991; accepted for publication 19 March 1991

Optimum composition and preparation conditions have been determined for improved proton-conductive 13"/13alumina precursors. The K ÷ content was decreased from 0.4 to 0.2 mole fraction, reducing the 13-aluminaphase fraction to < 0.15. Extendedtime ( _
I. In tro duct i o n

The high p r o t o n c o n d u c t i v i t y o f N H + - and H30 +13"-aluminas have generated much interest. These materials must be synthesized by ion-exchange o f precursor 13"-alumina ceramics that can withstand the necessary high t e m p e r a t u r e sintering. Na-13"-alumina is the preferred ceramic for sintering but it is destroyed during subsequent K+-ion-exchange expansion [ 1,2 ]. K+-13"-alumina cannot be sintered to a d e q u a t e density a n d f(13) (fraction o f 13-A1203 phase) as K ÷ favors the 13-A1203 phase. Research has e m p h a s i z e d the d e v e l o p m e n t o f a sinterable, mixedalkali ( N a / K ) c o m p o s i t i o n with low f(13). The ionic d i a m e t e r o f K ÷ (2.66 ,~) is similar to H 3 0 + (2.78 A ) . N a ÷ is m u c h smaller than either ( 1 . 9 0 / k ) . T h e level o f ion-exchange i n d u c e d stress in the sinter is inversely p r o p o r t i o n a l to its K + content but this K + p r o m o t e s f o r m a t i o n o f the 13-A1203 phase which has very low H3O+ c o n d u c t i v i t y ( 10-~l ( f l c m ) - l at 2 5 ° C for a single crystal as c o m p a r e d with the 13"-A1203phase ( 1 0 - 4 t o 10 -5 (~ c m ) -~ ) ). The original composition was thus designed with N a / K = 0 . 6 / 0 . 4 , (so-called " 6 N 3 " ) but had an f(13) ~ 0.4 [ 1 ]. 6N3 withstands the N a - K ion-exchange because the initial 40% K + content decreases the stresses ind u c e d during ion-exchange by ~ 40%. However, the resultant ceramics often failed during H 3 0 + - i o n - e x change. This p r o b l e m was investigated a n d residual

spinel (MgA1204) and NaA102 phases on grainb o u n d a r i e s were identified be the cause [3]. This p r o b l e m was solved by extending the sintering time [3]. The resulting ceramics satisfactorily withstand H30+-exchange in dilute acetic acid [2]. The rem a i n i n g 13-A1203 phase in the ceramics reduces their conductivity and the H 3 0 ÷ or NH~- ion-exchange rate. Steps were therefore taken to eliminate it. The level o f 13-A1203 in the sinter is controlled by its K ÷ content so it was reduced. The consequent increase o f ion-exchange induced stress must be a c c o m m o dated. This was achieved by increasing the temperature o f K + exchange to 1200°C to facilitate annealing. D a m a g e during the low-temperature K +H3O+ ion-exchange process was e l i m i n a t e d by grainb o u n d a r y structural i m p r o v e m e n t via extended sintering. The N a / ( K + N a ) c o m p o s i t i o n o f the precursor ceramics was increased from 0.6 to 0.8. The increased sintering time ensures complete reaction between any MgAI204 and 13-A1203 to give 13"-A1203. The 13"-A1203 phase stabilizing MgO content was also increased from 3% to 4%. This further lowers the final f(13). The resulting ceramics have an f(13) < 0.15 and a density > 3.15 g/cm 3 ( > 97% theoretical). This new precursor ceramic is t e r m e d " 8 N 4 " . The present p a p e r describes the characteristics o f 8N4 and compares its electrical a n d ion-exchange properties with those o f 6N3.

0167-2738/91/$ 03.50 © 1991 - Elsevier Science Publishers B.V. North-Holland )

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A. 7i~net al. / Proton-conductive fl" /fl-aluminas

2. The synthesis, phase content and microstrueture of 8N4

Nal.28K~.32Mgo.6oAl10.4oOlT. The powder synthesize

theoretical density is 3.25 g/cm 3) and f([~) <0.15. the microstructures of 6N3 and 8N4 sinters are shown in figs. l a and l b respectively. The preparation conditions and final properties of 6N3 and 8N4 are compared in table 1.

techniques are the same as for 6N3 [4]. During calcination (Na/K)A102 and MgA1204 as well as N a / K-[3"-AI203 are formed, i.e.;

3. Ion exchange of 8N4 ceramics

The

composition

of

8N4

is

( N a 2 0 + K20+AI203 + M g O )

3. I. Na +-K +-H30+ ion-exchange

(Na/K)AIO2 + MgAI204 + (Na/K)-[~"-AlzO3(f(~) ~ 0 ) .

( 1)

The high Na fraction and Mg content result in a powder containing almost no [3-A120 3 after calcination at 1280°C. The specific surface area is > 12 m 2 / g, which is higher than 6N3 (10-12 m2/g). This could be result of the increased Mg content raising the refractoriness of the calcine thus reducing preliminary sintering. The increased surface-area makes powder shaping difficult due to decreased pressure transmission on pressing. The latter induces green density gradients in the sample. Mold lubrication reduced density gradients so avoiding cracking of the final ceramic due to differential sintering shrinkage. The formation of MgA1204 means the Na/K-[3"A1203 phase formed on calcination contains insufficient Mg stabilizer and will transform to ]3-A1203at the higher temperatures employed for sintering ( 1610 °C). Two simultaneous reactions occur at the sintering temperature:

The precursor ceramics are ion-exchanged with K + at high temperature (1200~C) to expand the lattice to a size necessary to accommodate H30 + ions. The maximum potential stress induced in 8N4 is higher than in 6N3. The XRD-determined lattice parameters for 8N4 isomorphs are listed in table 2. The ex-

Na/K-[3"-A12 O3 ~Na/K-[~-AI203 + ( N a / K ) A I O 2 ,

(2)

( N a / K ) AIO2 + Na/K-I3-A12 03 + MgAI2 O4 Na/K-13"-A12 0 3 .

(3)

Reaction ( 3 ) is slower than reaction (2) so thef([3) increases during the early stages of sintering. The f(13) of a five minute sintered sample is ~0.3 and MgAI204 is detected in the sample [ 3 ]. Reaction (3) is complete after one hour sintering, i.e., f(]3) decreases to -<0.15 and the spinel-phase disappears. The sintering temperature for 8N4 is 10°C higher than for 6N3 because the higher Mg content increases the refractoriness of the system. The ceramics are crack free with density > 3.15 g/cm 3 (the

Fig. 1. SEM microstructuresof6N3 (a) and 8N4 (b) sinters.

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Table 1 Preparation and properties of 6N3 and 8N4. Composition XN~ 6N3 8N4

0.6 0.8

Powder MgO (wt%)

S a)

3 4

10-12 10-15

Sintering conditions f(fl)

(m2/g) <0.1 0

Sinters time (h)

f(fl)

d a)

(°C) 1610 1620

1 1

0.35-0.40 0 -0.15

0.96-0.99 0.97-0.99

T

a~ S is the specific surface area, d is the fraction of theoretical density.

Table 2 Lattice parameters of 8N41~"/1~-A1203isomorphs (A). Ion

Na+/K +

K+

Rb +

NH +

ao Coo,, Co~

5.620 _+0.001 33.64 _+0.01 22.35 _+0.01

5.624 _+0.001 34.05 _+0.01 22.70 _+0.01

5,630 _+0.001 34,34 _+0.01 22.95 _+0.01

5.631 _+0.001 34.34 _+0.02 22.90 _+0.02

pansion caused by K + exchange is higher then for 6N3 (Aao=0.002/~, Acorn,=0.31 A and Aco~=0.30 A, respectively [2] ). The higher volume strain develops higher macrostresses during K+-ion-exchange. Microstresses also develop and are determined by the strain differences in the Co and ao directions. The critical linear strain differential for microfracture initiation is [ 2 ]; j

,

(1)

where A I / I is the linear strain during ion-exchange and Co and ao the crystallographic directions. E is the Young's modulus o f the materials, Zbg the grain boundary fracture energy,/z the Poisson ratio and L the grain-size [2 ]. No matter how little 13-A1203 phase exists, the m a x i m u m microstress will be induced in it and microcracks will initiate there first. The m a x i m u m linear strain differential for 8N4 during K+-ion-ex change is, Aco~ Co~

Aao ~0.0149. ao

This figure is higher than for 6N3 (0.0134 [ 2 ] ). The 8N4 withstands pure KC1 vapor ion-exchange due to the high ion-exchange temperature ( 1200 ° C) which

ensures partial stress relief. The survival is also due to the stronger grain-boundaries o f 8N4 (tog). Failure during ion-exchange is due to propagation of microcracks under the macrostresses developed. This is a particular problem when the latter coincide with m a x i m u m microstresses [2]. The m a x i m u m macrostresses occur in 8N4 at the beginning of ion-exchange i.e., when the surface of the specimen is fully K+-exchanged and the interior has not started to ionexchange. This differential decreases as ion-exchange progresses. The induced tensile macrostress in the interior of the sample will induce propagation of existing microcracks therein. Microcracks initiate when the critical linear expansion differential (expressed by eq. ( 1 ) ) is reached. This critical value is proportional to the square-root o f the grain-boundary fracture energy (rbg). The extended sintering sequence reacts away the grain-boundary MgA1204 and ( N a / K ) A 1 0 2 so increasing the rgb. This in turn increases the critical value of the linear strain differential for crack initiation. This level may not be reached during K + ion-exchange, or, by the time it is reached, the macrostresses will have decreased to safe levels. K + - H 3 0 + ion-exchange is conducted with electric field assistance [ 1,5 ]. The ion-exchange rate o f 8N4 is 2 - 1 0 times that o f 6N3 due to the lower volume

A. Tan et al. / Proton-conductive ~" /~-alurninas

322

fraction of [ 3 - A 1 2 0 3 and the more perfect grainboundary structure. Achievable ion-exchange current levels increased to ~ 4 0 ~tA from ~ 10 ~tA at 80 ° C under 20 V, for a 1 mm thick, 1.1 cm 2 area pellet. The 8N4 is H30+-exchanged after ~ three weeks. 6N3 takes >__six weeks.

3.2. K+-Rb+-NH~ ion-exchange The lattice parameters of Rb +-8N4 are listed in table 2. The maximum strain differential for the K +Rb + ion exchange is;

-]- c~o

-)- .o=0"0099"

tiated microcracks will propagate and break the sample eventually. Rb+-NH~- ion-exchange of 6N3 is very slow ( > 20 days). The high content of [~-A1203 is difficult to NH~- ion-exchange. NHg ion-exchange of 8N4 is much faster ( < 10 days). Some samples completely ion-exchanged in two days. This variation of ion-exchange rate is not understood. It may be associated with the microstructure, grain-boundary structure a n d / o r the residual ceramic composition following the Na-K-Rb ion-exchange processes. Conductivity measurements could not explain the variable NH + ion-exchange rates.

4. The ae-conduetivity of the 6N3 and 8N4 analogues

This is smaller than for the Na +-K + ion-exchange. The microstresses induced by the initial Na-K ionexchange cannot be removed completely if no microcracks are created. Microcracks will initiate when the summed stresses accumulate to the requisite value. A smaller difference between the K +- and Rb+-8N4 lattice parameters lowers the maximum macrostress value so microcracks, if present, will not propagate. The material thus withstands K+-Rb + ion-exchange. Direct K+-NH + ion-exchange is not possible. 8N4 pellets were destroyed after 10 days ion-exchange due to the necessary low temperature ( ~ 200 ° C). Stresses are not relieved at such low temperatures. Microcracks initiate when the accumulated microstress reaches the critical value. Slow ion-exchange rates maintain high macrostresses in the sample and ini-

The ac-impedance-derived conductivities of the 8N4 and 6N3 [3"/13-A1203isomorphs are listed in table 3. 8N4 has improved grain and grain-boundary conductivity. The increase of grain conductivity is due to increased [3"-A1203 phase levels. The conductivity increase of the grain-boundaries is more significant (1.4 times to one order of magnitude). This improvement is attributed to the improvement of grain-boundary structure during the extended sintering procedure. The latter assures second-phasefree, less-distorted grain-boundary structures and thus lower resistance. The lower level of ~ - A I 2 0 3 phase will also increase the ionic conductivity of the grainboundaries as fewer [~-[3- and [3-[~"-A1203 resistive boundaries exist [ 6 ]. The conductivity - 1/ T curves for NH + - and H30+-8N4 and 6N3 are shown in figs.

Table 3 C o n d u c l i v i t i e s o f 6 N 3 a n d 8 N 4 [5"/[~-A1203 i s o m o r p h s ( f l - J c m - ~ ). Ion

Na÷/K ÷ K+ N H 4+ H3 O +

T (°C)

25 200 25 200 25 200 25 20O

Grain-boundaries

Grain 6N3

8N4

6N3

8N4

8.1 × 10 -5 3.7)<10-3 1.4X 10 -5

1 . 4 × 10 -3

3 . 0 × 10 -3

5.8×10 3 2 . 0 × 10 - 2

1 , 1 × 10-~ 4 , 9 X 10 -2

3 . 0 X 10 -1

4 , 5 X 10 -1

1.8X 10 -3

8 . 3 X 10 - 5

1.3X 10 - 4

2 . 7 X 10 -3

3 . 5 × 10 -3

3.9)< 10 -6 6.3)< 10 - 4

5.4)< 10 - 6 1.1 )< 10 -3

5.5X 2.4× 1.5)< 3.2)<

10 -6 10 - 4 10 - 6 10 -5

6 . 7 × 10 -4

2.1X 10-2 7.6×10 5 2 . 7 X 10 -3 1.1X 10 -5 5.4)< 10 -4

2.6)< 10 -6 2.3)< 10 -4

A. Tan et aL / Proton-conductive fl" /#-aluminas Temperoture(°C) 200 1O0 I

-2.5

~

l

20 I

-3.5 -

7 E vb

50

I

4

-4.5 -

-5.5

1.5

2.b

3.b

3.5

1000/T(K)

Fig. 2. Conductivity versus 1000/Tofammonium 6N3 and 8N4. Temperoture(°C) 200 1O0 I

-2.5

I

~" -3.5

50 I

20 I

\ Grain

T

o

~ -4.5 o b

-5.5

1.5

2.b

2.~ IO00/T(K)

curve in fig. 3 tracks the grain conductivity versus temperature for H30+-6N3). Unlike other 13"/13A1203 isomorphs, grain-conduction has a higher activation energy than grain-boundary conduction in H 3 0 +-13''/13"A1203 [ 6 ].Clearly the bulk conductivity of H30+-8N4 is controlled by a microstructural entity similar to the grains, having switched thereto from the boundary-control of 6N3. This must be due to the improvement of the boundary structure. The equality of the AE for NH~--6N3 and 8N4 conduction suggests the same process is rate-controlling ionic motion in both materials but the mobility of the ions increases. It is puzzling that the behaviour of NH~--8N4 versus N H 2 - 6 N 3 does not mirror the H30 + analogues. It could be that the motion of the mobile species (H + in NH~- -13"-A1203 [ 7 ] ) is controlled by the same microstructural entity albeit improved (less restrictive) in NH~--8N4. the mobile species in H30+-13 "A1203 is H3 O + not H ÷ so it is not surprising it is sensitive to microstructural improvement.

5. Summary

\\

T

-6.5

323

3.b

3.5

Fig. 3. Conductivity versus 1000/T ofhydronium 6N3 and 8N4.

2 and 3 respectively. The improvement is obvious but the rate-controlling mechanism of H30 + conduction appears to change from 6N3 to 8N4. It could be the decreased level of the highly resistive 13-A1203 phase in the H30+-8N4 results in an increased activation energy ( z ~ k E H 3 0 + . 6 N 3 = 0.22 eV, LkEH3o+_8N4=0.35 e V ) but this is unlikely. It is tempting to associate the change with the cleaner grain-boundaries and better crystal-lattice fit between adjoining grains. The rate-controlling conduction mechanism of H3 O + in 8N4 is more similar to that in grains than is the case for 6N3 (the dotted

The development of improved Na-K-13"/13-A1203 precursor ceramics for H 3 0 +- and NH~" -13"/13-A1203 electrolyte synthesis is described. The new 13"/13A1203 material (8N4) contains less 1 3 - A 1 2 0 3 ( -~ 0.15 ) and improves grain boundary structure. These properties led to improvement of the ion-exchange rates and the conductivity of the precursor ceramics. The lower K + content of the precursor ceramic ensures a lower f(13) and the concommitant disadvantage of higher stress induction during K+-ion-exchange is handled by increasing the alkali exchange temperature and the stronger grain-boundaries. There is evidence that the ionic-motion rate-controlling microstructural entity change from grain boundaries to grains in the H30+-8N4 whereas the ionic motion in the N H ~ - 8 N 4 is controlled by the same process as in NH~--6N3.

References [ 1 ] P.S. Nicholson, M.Z.A. Munshi, G. Singh, M. Sayer and M.F. Bell, Solid State lonics 18/19 (1986) 699.

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A. Tan et al. / Proton-conductive fl" ~,O-aluminas

[ 2 ] A. Tan and P.S. Nicholson, Solid State lonics 26 ( 1988 ) 217. [3] A. Tan and P.S. Nicholson, Solid State Ionics 37 (1989) 51. [4] P.S. Nicholson, M. Nagai, K. Yamashita M. Sayer and M.F. Bell, Solid State lonics 15 (1985) 317.

[ 5 ] A. Tan, C.K. Kuo and P.S. Nicholson, A study of hydroniumexchange of polycrystalline Na-K-I~"/[~-AlzO3, Solid State Ionics, submitted. [6] A. Tan, C.K, Kuo and P.S. Nicholson, Solid State Ionics 45 (1991) 137. [7] J.O. Thomas, Acta Cryst. B 39 (1983) 227.