J O U R N A L O F M A T E R I A L S S C I E N C E L E T T E R S 2 2, 2 0 0 3, 1779 – 1781

Effect of nitridation on the properties of Na–Zn–P–O glasses P. Y. SHIH Department of Ceramic and Materials Engineering, National Lien Ho Institute of Technology, Miaoli 36012, Taiwan, People’s Republic of China E-mail: [email protected]

C 2003 Kluwer Academic Publishers 0261–8028 

on polished glass samples. Glass powders were used for the differential thermal analysis (DTA) at a heating rate of 10 ◦ C/min to determine the glass transition temperature (Tg ). The relative chemical durability was estimated by measuring the dissolution rate of the polished glasses, which were immersed in deionized water for 24 h at 30 ◦ C. The nitrogen content in the glasses was analyzed by an inert-gas fusion method, using an oxygen-nitrogen analyzer. X-ray photoelectron spectroscopy (XPS) spectra of N 1s core levels were recorded by a V. G. Scientific EscaLab 210 XPS spectrometer. The X-ray source was Mg Kα (hν = 1253.6 eV). The nitrogen contents of the nitrided glasses with various remelting times are shown in Fig. 1. The nitrogen content of the glasses increases with increasing remelting time, reaching a maximum value of 1.81wt% (2.9 at.%) for a remelting time of 5 h. The dissolution rate and concentration of nitrogen in the glass network are affected by glass composition and the remelting conditions. It was suggested that higher temperatures significantly accelerate nitrogen dissolution [11, 12, 26]. Alkali phosphorus oxynitride (Na P O N and Li P O N) and sodium alkalineearth metaphosphorus oxynitride glasses have been studied [10, 25, 26]. Lithium phosphorus oxynitride glasses that were prepared at 700 ◦ C for 48 h had a maximum nitrogen content of 9.2 wt%, whereas those glasses that were prepared at 800 ◦ C had a value of 12.6 wt% [12]. The nitrogen solubility in sodium alkaline-earth metaphosphorus oxynitride glasses decreased with increasing concentration of alkaline-earth oxides [26]. It was suggested that the divalent cations 2 Nitrogen content (wt %)

Phosphate glasses generally have a low melting temperature, a low glass transition temperature (Tg ) and a high thermal expansion coefficient (α) [1–6]. These characteristics make them potential candidates for low temperature applications, such as the molding of optical elements [7] and sealing to high expansion metals [8]. However, the applications of phosphate glasses are greatly limited in practice by their poor chemical durability [6, 9]. It has been shown that the chemical durability of phosphate glasses could be significantly improved by nitridation and/or addition of appropriate additives, due to the increase of crosslink density and bonding strength in the glass network [9–18]. The effect of nitridation on the glass properties results in the two-coordinated oxygen atom being replaced by the three-coordinated nitrogen atom [19–22]. Either doping the melt with metal nitrides or remelting the glass in an ammonia atmosphere can lead to the incorporation of nitrogen atoms into the glass network [23–25]. Higher nitrogen contents can be obtained by the remelting method than by doping with metal nitrides [10, 24]. It has been shown that the nitrogen content of nitrided glasses would increase with increased remelting temperature and remelting time, but this may make melt volatilization more likely [10–12]. The purpose of this work is to improve the chemical durability of sodium-zinc phosphate (Na Zn P O) glasses by remelting the base glasses at a lower temperature and a shorter holding time under an ammonia (NH3 ) atmosphere. A base glass with composition 50P2 O5 –20Na2 O– 30ZnO (mol%) was investigated in this study. The base glasses were prepared from mixtures of reagent grade sodium dihydrogen phosphate (NaH2 PO4 ), zinc oxide (ZnO), and ammonium dihydrogen phosphate (NH4 H2 PO4 ). Well-mixed powders were melted in an alumina crucible in air at 850 ◦ C for 1 h. The melts were quenched and crushed to frits. Then the crushed base glasses were nitrided in alumina boat crucibles under a flowing anhydrous ammonia atmosphere at 600 ◦ C for different durations, from 0.5–5 h. The ammonia flow rate was kept at 300 cm3 /min. After selected periods of time, the glass samples were removed and annealed in a preheated furnace under an air atmosphere. Then the oxynitride glass samples were cooled overnight to room temperature. Density measurements were carried out at room temperature, using the Archimedes method with deionized water as the immersion fluid. A micro-Vickers tester was used to measure the hardness with a 300 g load

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Remelting time (h) Figure 1 Nitrogen content of sodium-zinc phosphate glasses nitrided at 600 ◦ C for various remelting times.

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=N− −NH−

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Figure 3 Dependence of density and hardness on remelting time for the nitrided sodium-zinc phosphate glasses.

Figure 2 N 1s XPS spectrum of the sodium-zinc-phosphate glass nitrided at 600 ◦ C for 5 h.

will reduce the diffusion of NH3 in the phosphate glass melts by filling the interstitial sites in the PO4 network and lead to decreases in the nitrogen content of the oxynitride glasses [26]. The nitrogen content of 30Na2 O–20BaO–50P2 O5 remelted at 650 ◦ C for 5 h had a value of 2.05 wt% [27]. The nitrogen content of the studied Na Zn P O glasses is lower than that of alkali phosphorus oxynitride glasses but is similar to the value of 30Na2 O–20BaO–50P2 O5 glasses. We suggest that the Zn2+ cations in the studied glasses may have a similar effect to that of the Ba2+ ion, which interacts with PO4 chains by forming ionic cross-linkages between the non-bridging oxygen (NBO) that belong to different chains and reducing the diffusion of NH3 in the phosphate melts. Fig. 2 depicts the N 1s XPS spectrum of 50P2 O5 – 20Na2 O–30ZnO glass that had been nitrided at 600 ◦ C for 5 h. The spectrum can be decomposed into three components that were located at about 397.8, 399.3 and 400.8 eV. These components are attributed to three different bonding states of nitrogen atoms, N , N< and NH , respectively [11, 28]. The content of the NH group, which can replace the terminal OH ions in the glass structure, is relatively small. Previous work found that most of the incorporated nitrogen was present as nitride ions, N3+ , ( N< and N ), which will replace the bridging oxygen (BO, O ) and nonbridging doubly bonded oxygen ( O) [21, 29]. These reactions can be represented as follows:

The density and hardness of the studied glasses are shown in Fig. 3. Density and hardness both increase with increasing remelting time. This indicates that nitriding the base glass will strengthen the glass network and make it more compact. The identity of the modified cation and the phosphate chain length affect the density of the glasses, with the higher field strength cations and shorter chain lengths resulting in higher densities. The relative content of modified cations (Na+ and Zn2+ ) is unchanged. Therefore, the increasing density of the glasses is attributed to the shortened phosphate chain length and the decreasing relative content of BO as nitrogen is successively incorporated into the glass network in the nitriding process. The dependence of the dissolution rate and Tg on the remelting time of the glasses is illustrated in Fig. 4. The dissolution rate decreases but Tg of the glasses increases with increasing remelting time. For phosphate glasses, the value of Tg increases with increasing phosphate chain length and cross-link density of the structure. As discussed above, phosphate chain length decreases with increasing nitrogen content. Therefore, the increase of Tg with nitrogen content is due to an increase in the cross-link density of the structure. When nitrogen incorporates into the glass network, the functional groups ( NH , N and N<) replace the terminating hydroxyl ion ( OH), bridging oxygen ( O ) and nonbridging oxygen ( O), which leads to enhanced cross-linking of the glass network. The decrease in the dissolution rate also indicates a stronger and higher cross-link density in glass network.

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3( O )(glass) + 2NH3 → 2( N<)(glass) + 3H2 O ↑ (1) Tg (oC)

(2)

where two N< can replace three BOs, and two N can replace one BO and two doubly bonded oxygen. The substitution of triply coordinated nitrogen ( N<) for doubly coordinated oxygen ( O , and the doubly coordinated nitrogen ( N ) for singly coordinated oxygen ( O) will render a more highly crosslinked network. 1780

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( O )(glass) + 2( O)(glass) + 2NH3 → 2( N )(glass) + 3H2 O ↑

-4 Tg Dissolution rate

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404 402 400 398 396 394 Binding energy (eV)

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Figure 4 Glass transition temperature (Tg ) and dissolution rate (D.R.) as a function of remelting time for nitrided sodium-zinc phosphate glasses.

Metaphosphate glasses “(M2 O + MO)/P2 O5 = K = 1” consist of phosphate chains, which are bound together by cations between them. The chemical durability of such metaphosphate glasses is affected by the following factors: (1) the strength of ionic bonding between cations and phosphate chains, (2) the attractive force between cations and the surrounding H2 O, and (3) whether the H2 O molecules can enter easily into the glass structure between phosphate chains. The functional groups N and N<, which link between phosphate chains, form in the nitrided glasses network. Therefore, the incorporation of nitrogen not only strengthens the glass network but also retards the entrance of H2 O, and leads to improved chemical durability of the glasses. Acknowledgment The author is grateful to the National Science Council of the Republic of China for financial support, under the grant NSC 88-2216-E-239-002. References 1. P . Y . S H I H , S . W . Y U N G , C . Y . C H E N , H . S . L I U and T . S . C H I N , Mater. Chem. Phys. 50 (1997) 63. 2. P . A . T I C K , Phys. Chem. Glasses 25(6) (1984) 149. 3. L . M . S A N F O R D and P . A . T I C K , U.S. Pat. 4314031, Feb. 2 (1982). 4. Y . H E and D . E . D A Y , Glass Technol. 33(6) (1992) 214. 5. H . S . L I U , P . Y . S H I H and T . S . C H I N , Phys. Chem. Glasses 37(6) (1996) 227. 6. C . M . S H A W and J . E . S H E L B Y , ibid. 29(2) (1988) 49. 7. X . J . H U and D . E . D A Y , ibid. 31(5) (1990) 183.

8. Y . B . P E N G and D . E . D A Y , Glass Technol. 32(5) (1991) 166. 9. B . C . B U N K E R , G . W . A R N O L D , M . R A J A R A M and D . E . D A Y , J. Amer. Ceram. Soc. 70(6) (1987) 425. 10. M . R . R E I D I M E Y E R , M . R A J A R A M and D . E . D A Y , J. Non-Cryst. Solids 85 (1986) 186. 11. H . Y U N G , P . Y . S H I H , H . S . L I U and T . S . C H I N , J. Amer. Ceram. Soc. 80(9) (1997) 2213. 12. R . W . L A R S O N and D . E . D A Y , J. Non-Cryst. Solids 88 (1986) 97. 13. C . M . S H A W and J . E . S H E L B Y , Phys. Chem. Glasses 29(3) (1988) 86. 14. I . W . D O N A L D , J. Mater. Sci. 28 (1993) 2841. 15. C . M . S H A W and J . E . S H E L B Y , J. Amer. Ceram. Soc. 71(5) (1988) C-252. 16. R . K . B R O W , C . M . A R E N S , X . Y U and D . E . D A Y , 35(3) (1994) 132. 17. R . K . B R O W , J. Amer. Ceram. Soc. 76(4) (1993) 913. 18. R . K . B R O W , R . J . K I R K P A T R I C K and G . L . T U R N E R , ibid. 76(4) (1993) 919. 19. H . O . M U L F I N G E R , ibid. 49(9) (1966) 462. 20. R . K . B R O W and C . G . P A N T A N O , ibid. 67(4] (1984) C-72. 21. B . C . B U N K E R and D . R . T A L L A N T , ibid. 70(9) (1987) 675. 22. M . R . R E I D I M E Y E R and D . E . D A Y , J. Mater. Res. 6(8) (1991) 1757. 23. Y . B . P E N G and D . E . D A Y , Glass Technol. 32(5) (1991) 166. 24. M . R . R E I D I M E Y E R and D . E . D A Y , J. Amer. Ceram. Soc. 68(8) (1985) C-188. 25. R . M A R C H A N D , J. Non-Cryst. Solids 56 (1983) 173. 26. M . R A J A R A M and D . E . D A Y , J. Amer. Ceram. Soc. 70(4) (1987) 203. 27. Idem., J. Mater. Sci. 24 (1989) 573. 28. R . K . B R O W , M . R . R E I D I M E Y E R and D . E . D A Y , J. Non-Cryst. Solids 99 (1988) 178. 29. D . E . D A Y , ibid. 112 (1989) 7.

Received 26 February and accepted 9 July 2003

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