Grain Boundary Conductivity in Crystalline LiTi2(PO4)3

by Jeff Wolfenstine 

ARL-TN-0313 April 2008 

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Army Research Laboratory 

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ARL-TN-0313 April 2008 

Grain Boundary Conductivity in Crystalline LiTi2(PO4)3

Jeff Wolfenstine 

Sensors and Electron Devices Directorate, ARL 

Approved for public release; distribution unlimited. 
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Grain Boundary Conductivity in Crystalline LiTi2(PO4)3

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14. ABSTRACT 

This Technical Note suggests that the low values of Li-ion grain boundary conductivity in crystalline LiTi2(PO4)3, compared 

to bulk Li-ion conductivity, are the result of a space charge region associated with grain boundaries and not a result of an 

amorphous phase along grain boundaries—as was previously suggested. This conclusion is based on microstructural 

evidence, a comparison of the activation energy for grain boundary conductivity, and recent results on doped polycrystalline 

oxygen-ion conductors containing no amorphous phases along grain boundaries. 

15. SUBJECT TERMS

Grain boundary, ionic conductivity, space charge, amorphous phase 

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Jeff Wolfenstine 

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Contents 

List of Figures iv

Acknowledgments v

References 4 

Distribution List 5 
iv 

List of Figures 

Figure 1. Schematic diagram of grain boundaries showing: a) amorphous phase and b) 

grain to grain contact (13)..........................................................................................................3 
v 

Acknowledgments 

The author would like to acknowledge support of the U. S. Army Research Laboratory. 
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1 

Recently, there has been a renewed interest in the development of high energy Li-Air batteries. 

One configuration involves the use of a Li anode in a non-aqueous electrolyte, which is separated 

from an aqueous electrolyte containing the air cathode by a solid state Li-ion conducting 

membrane. Several solid state polycrystalline Li-ion conductors, based on pervoskite (1,2), 

garnet (3,4) and NASICON (Na super ion conductor) (5 through 12) structures, are under 

consideration as possible membrane materials. One of most widely investigated crystalline Liion conducting membrane materials, based on the NASICON structure, is LiTi2(PO4)3 (5 through 

10). In order to sinter crystalline LiTi2(PO4)3 to the high relative densities required for use as a 

membrane and increase Li-ion conductivity, two approaches have been undertaken (5 through 

10). The first it to use a doped material, LiMxTi2–x(PO4)3 (where M=Al, Sc, Y and La) 

(5,6,9,10). The second is to use LiTi2(PO4)3 containing a small amount of Li2O or Li3PO4 or 

Li3BO3 (5,7,8). Previous investigations have suggested the total Li-ion conductivity, based on 

analysis of ac impedance data, of M-doped LiTi2(PO4)3 (where M=Al, Sc, Y and La) and 

LiTi2(PO4)3 containing a small amount of Li2O or Li3PO4 or Li3BO3 was controlled by Li-ion 

grain boundary conductivity, which is about 1 to 2 orders of magnitude lower compared to Li-ion 

bulk conductivity (5 through 10). It has been suggested that for both polycrystalline M-doped 

LiTi2(PO4)3 (where M=Al, Sc, Y and La) and LiTi2(PO4)3 containing a small amount of Li2O or 

Li3PO4 or Li3BO3, that both approaches lead to the formation of a continous amorphous film 

around the grains (5 through 10). It is the transport of Li-ions through this amorphous film 

which controls the sintering rate (i.e., densification) and grain boundary Li-ion conductivity, and 

hence, total Li-ion conductivity of the material (5 through 10). 

This Technical Note suggests an alternative interpretation of what controls Li-ion grain boundary 

conductivity in polycrystalline M-doped LiTi2(PO4)3 (where M=Al, Sc, Y and La) and 

LiTi2(PO4)3 containing a small amount of Li2O or Li3PO4 or Li3BO3, based on the existing 

experimental data for these materials and a comparison to recent results for polycrystalline 

oxygen-ion conducting membranes. This result has implications not only for increasing the total 

Li-ion conductivity of NASICON based Li-ion conductors but, also for polycrystalline Li-ion 

conductors based on pervoskite and garnet structures where the total Li-ion conductivity is also 

controlled by grain boundary Li-ion conductivity (1 through 5). 

For the case of polycrystalline M-doped LiTi2(PO4)3 (where M=Al, Sc, Y and La) and 

LiTi2(PO4)3 containing a small amount of Li2O or Li3PO4 or Li3BO3, it is postulated that the 

reason for the low grain boundary Li-ion conductivity is not due to Li-ion transport through the 

grain boundary amorphous phase as previously suggested but, instead is through grain to grain 

contact where the resistance is determined by a space charge region associated with the grain 

boundary. The presence of the amorphous phase serves only to reduce the amount of grain/grain 

contacts, as shown in figure 1 (13). Evidence to support the above claim is as follows: 

Firstly, no conclusive microstructural evidence for an amorphous phase along grain boundaries 

in polycrystalline M-doped LiTi2(PO4)3 (where M=Al, Sc, Y and La) and LiTi2(PO4)3 containing 

a small amount of Li2O or Li3PO4 or Li3BO3 materials exists (5,7,9). For example, scanning 
2 

electron micrographs do not reveal any rounding of grain corners at multiple grain junctions 

(7,9), which is typical for material containing an amorphous boundary phase (14 through 16). 

Furthermore, no transmission electron micrographs showing evidence of an amorphous phase 

along the grain boundaries are presented. 

Secondly, if a grain boundary phase controls grain boundary conductivity, then the activation 

energy for grain boundary conductivity should be different for chemically different boundary 

phases (13,17 through 19,20). This is in contrast to the experimental results where the activation 

for grain boundary conduction in the different M-doped LiTi2(PO4)3 (where M=Al, Sc, Y and 

La) materials and LiTi2(PO4)3 containing a small amount different additives such as; Li2O or 

Li3PO4 or Li3BO3 all have nearly the same value (∼0.36 eV [5]). 

Thirdly, recent results for doped polycrystalline CeO2 and ZrO2 oxygen-ion conductors have 

shown that the oxygen-ion grain boundary conductivity is 2 to 3 orders of magnitude lower than 

for bulk oxygen-ion conductivity (13,19), in agreement with the results for M-doped LiTi2(PO4)3

(where M=Al, Sc, Y and La) materials and LiTi2(PO4)3 containing a small amount of Li2O or 

Li3PO4 or Li3BO3 materials. For the case of these oxygen-ion conductors, detailed electron 

microscopy revealed clean grain boundaries (no amorphous phases along the boundaries). Any 

amorphous phase was located at triple junctions and not along grain boundaries. Thus, the 

presence of an amorphous phase along the grain boundaries is not required to account for low 

grain boundary ionic conductivity compared to bulk ionic conductivity. Maier et al., (13,17 

through 19) have suggested that the major reason for the low grain boundary ionic conductivity 

in these materials is a result a space charge effect associated with grain boundaries compared to 

the bulk, where no space charge exits. The presence an amorphous phase serves only to reduce 

the fraction of grain/grain contacts (figure 1). The space charge model has successfully 

predicted the ratio of grain boundary conductivity/bulk conductivity, activation energy for grain 

boundary diffusion and the effect of aliovalent cation dopants on grain boundary conductivity in 

the doped polycrystalline CeO2 and ZrO2 oxygen-ion conductors (13,19). 
3 

Figure 1. Schematic diagram of grain boundaries showing: a) amorphous 

phase and b) grain to grain contact (13). 

In conclusion, from microstructural evidence similar values of the activation for grain boundary 

conductivity of the LiTi2(PO4)3 material with different dopants and second phase materials, and 

from recent results on doped polycrystalline oxygen-ion conductors, where no amorphous phases 

along grain boundaries were observed, whose lower grain boundary oxygen-ion conductivity to 

bulk oxygen-ion conductivity can be explained by a space charge model, it is highly likely that 

the low value of Li-ion grain boundary conductivity compared to bulk Li-ion conductivity, 

reported previously in crystalline M-doped LiTi2(PO4)3 (where M=Al, Sc, Y and La) and 

LiTi2(PO4)3 containing a small amount of Li2O or Li3PO4 or Li3BO3 materials, is through direct 

grain to grain contact where the resistance is controlled by a space charge region associated with 

the grain boundaries and not a result of an amorphous phase along grain boundaries, as 

previously suggested. The presence of any amorphous phase serves only to reduce the amount of 

grain/grain contacts. To conclusively confirm this suggestion, a very detailed investigation 

relating microstructure to conductivity in high density pure polycrystalline LiTi2(PO4)3 and 

LiTi2(PO4)3 containing a second phase is required. The results of this Technical Note have 

implications not only for increasing the total Li-ion conductivity of NASICON based Li-ion 

conductors but, also for polycrystalline Li-ion conductors based on pervoskite and garnet 

structures, where their total Li-ion conductivity is also controlled by grain boundary Li-ion 

conductivity. 

 A 

 B 

 A 
4 

References 

1. Thangadurai, V.; Weppner, W. Solid State Ionics 2006, 12, 81. 

2. Inaguma, Y.; Chen, L. Q.; Itoh, M.; Nakamura, T.; Uchida, T.; Ikuta, H.; Wakihara, M. 

Solid State Commun. 1993, 86, 689. 

3. Thangadurai, V.; Schwenzel, S.; Weppner, W. Ionics 2005, 11, 11. 

4. Thangadurai, V.; Weppner, W. J. Am. Ceram. Soc. 2005, 88, 411. 

5. Adachi, G.; Imanaka, N.; Aono, H. Adv. Mater. 1996, 8, 127 

6. Aono, H.; Sugimoto, E.; Sadaoka,Y.; Imanaka, N.; Adachi, G. Solid State Ionics 1990, 40–

41, 38. 

7. Aono, H.; Sugimoto, E.; Sadaoka,Y.; Imanaka, N.; Adachi, G. Solid State Ionics 1991, 47, 

257. 

8. Aono, H.; Sugimoto, E.; Sadaoka,Y.; Imanaka, N.; Adachi, G. J. Electrochem. Soc. 1993, 

140, 1827. 

9. Aono, H.; Sugimoto, E.; Sadaoka,Y.; Imanaka, N.; Adachi, G. J. Electrochem. Soc. 1990, 

137, 1023. 

10. Aono, H.; Sugimoto, E.; Sadaoka,Y.; Imanaka, N.; Adachi, G. Chem Lett. 1990, 1825. 

11. Li, S. C.; Cai, J. Y.; Lin, Z. X. Solid State Ionics 1988, 28–30, 265. 

12. Delmas, C.; Nadiri, A.; Soubeyroux, J. L. Solid State Ionics 1988, 28–30, 419. 

13. Guo, X.; Maier, J. J. Electrochem. Soc. 2001, 148, E121. 

14. Barsum, M. V. Fundamentals of Ceramics; The McGraw-Hill Companies, Inc. New York, 

1997. 

15. Chiang, Y. M.; Birnie III, D.; Kingery, W. D. Physical Ceramics; Wiley, New York, 1997. 

16. Kingery, W.D.; Bowen, H.K.; Uhlmann, D.R. Introduction to Ceramics (second ed); Wiley, 

New York, 1976. 

17. Maier, J. Solid State Ionics 2000, 131, 13. 

18. Guo, X.; Sigle, W.; Fleig, J.; Maier, J. Solid State Ionics 2002, 154–155, 555. 

19. Guo, X.; Sigle, W.; Maier, J. J. Am. Ceram. Soc. 2003, 86, 77. 

20. Badwal, S.P. S. Solid State Ionics 1995, 76, 67. 
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