Effects of ball milling time on microstructure evolution and optical transparency of Nd:YAG ceramics

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CERAMICS INTERNATIONAL

Ceramics International 40 (2014) 9841–9851 www.elsevier.com/locate/ceramint

Effects of ball milling time on microstructure evolution and optical transparency of Nd:YAG ceramics Jun Liua, Li Lina,b, Jiang Lib,n, Jing Liua,b, Yong Yuana,b, Maxim Ivanovc, Min Chenb, Binglong Liub, Lin Geb, Tengfei Xieb,d, Huamin Koub, Yun Shib, Yubai Panb,n, Jingkun Guob a School of Materials Science and Engineering, Jiangsu University, 301 Xuefu Road, Zhenjiang 212013, China Key Laboratory of Transparent and Opto-functional Advanced Inorganic Materials, Shanghai Institute of Ceramics, Chinese Academy of Sciences, 1295 Dingxi Road, Shanghai 200050, China c Institute of Electrophysics, Ural branch of Russian Academy of Sciences, 106 Amundsena street, Ekaterinburg, Russia d Key Laboratory of Resource Chemistry of Education Ministry, Department of Chemistry, Shanghai Normal University, 100 Guilin Road, Shanghai 200234, China b

Received 30 January 2014; received in revised form 19 February 2014; accepted 20 February 2014 Available online 26 February 2014

Abstract Polycrystalline Nd:YAG ceramics were fabricated by the solid-state reaction and vacuum sintering method using Y2O3, α-Al2O3 and Nd2O3 as starting powders. These powders were mixed in ethanol doped with MgO and TEOS and ball milled for different time periods. The samples were sintered from 1500 1C to 1760 1C for 0.5–20 h. Effects of ball milling time on the particle size of powder mixtures as well as ceramics densification process, microstructure evolution and optical transparency of the as-prepared Nd:YAG ceramics were mainly investigated. It was demonstrated that coarse powders can be ground into fine particles. Porosities of Nd:YAG ceramics sintered at 1600 1C and 1760 1C decrease with the increase of ball milling time up to 12 h and are kept almost unchanged with further increasing the time. Contamination of the powders with impurities, which may not be observed in microstructures, when the ball milling time is overlong, and the very few residual micro-pores, will cause an infinitesimal decrease in transmittance. The grain growth kinetics of Nd:YAG ceramics fabricated from the optimal ball milling process was also studied in this paper. The highest in-line transmittance of 83% at 1064 nm was obtained by sintering the sample at 1760 1C/3 h from the powder mixtures ball milled for 12 h. & 2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

Keywords: Ball milling time; Densification trajectory; Microstructure evolution; Nd:YAG ceramics; Optical transparency

1. Introduction Laser materials have been widely used in various fields such as metal processing, micromachining, medical applications, environmental instrumentation measurements and optical transmission systems [1–3]. One of the most important laser material is Nd:YAG single crystal due to its excellent optical and thermal properties, which is usually produced by the Czochralski (CZ) method [2]. However, it is hard to grow large sized Nd:YAG crystals for the limitation of crystal growth n

Corresponding authors. Tel.: þ86 21 52412816; fax: þ 86 21 52413903. E-mail addresses: [email protected] (J. Li), [email protected] (Y. Pan). http://dx.doi.org/10.1016/j.ceramint.2014.02.076 0272-8842 & 2014 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

technique. Moreover, owing to its low effective segregation coefficient of neodymium in YAG single crystal, it is difficult to dope neodymium homogeneously with high doping concentration [4]. Ikesue et al. first realized laser oscillation in Nd: YAG ceramics, which was obtained by the solid state reactive sintering process in 1995 [5]. Since then, the fabrication of transparent YAG ceramics has got a significant improvement. Lu et al. and Ikesue demonstrated the laser properties of Nd: YAG transparent ceramics which were nearly equivalent or even superior to those of high-quality Nd:YAG single crystals [4,6,7]. Recently Nd:YAG transparent ceramic has gained considerable attention as a substitute for Nd:YAG single crystal due to its low cost, short preparation period, high doping concentration, excellent laser performance, ease of

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mass production [1,6,8] and so on. Generally, there are two typical methods to fabricate YAG ceramics: one is solid-state reactive sintering of primary oxides powders mixed in stoichiometric proportions [9–12], and the other is sintering of YAG powders obtained by the wet-chemical method [13–19]. Solid-state reaction and vacuum sintering method is a comparatively simple process to fabricate YAG ceramics, and it is easy to achieve a series of composition by changing the powder amounts. Lots of the papers have focused on starting powders selection or synthesis [20–22], forming method [23–26] and ceramic sintering process [9,27,28], as well as analysis of doping mechanism [29,30], function of sintering aids [31–35] and optical properties [36–39]. However, effects of ball milling process on the microstructure and optical transmittance of Nd:YAG ceramics fabricated by solid-state reactive sintering were seldom reported. In this work, high purity commercial powders were used as raw materials, and highly transparent Nd:YAG ceramics were obtained by the solid-state reaction and vacuum sintering method. Effects of different ball milling time on micromorphology of powder mixtures, densification process, microstructure evolution and optical transparency of the obtained ceramics were mainly investigated. The grain growth kinetics of Nd:YAG ceramics fabricated by the optimized ball milling process was also studied in this paper.

(tetraethyl orthosilicate, 4 99.99%, Alfa Aesar, USA) as sintering aids. Then the mixed powders were milled on a planetary ball mill for 2–20 h with 10 mm diameter alumina balls in ethanol. The disc and bottle rotation speeds were 130 rpm and  260 rpm, respectively. The mass ratio of balls and powders was 3:1 and the solid loading of the ball milled slurry was about 1.8 g/ml. After dried and sieved through a 200-mesh screen, the powder mixtures were calcined at 800 1C to remove organic components and then uniaxially dry-pressed at 60 MPa to form disks with 20 mm diameter, and cold isostatically pressed at 250 MPa. The obtained green pellets were vacuum-sintered at 1760 1C for 3 h ( r 10  3 Pa). After sintering, the specimens were annealed at 1450 1C for 20 h in air atmosphere to get rid of the oxygen vacancies. Finally the specimens were mirror-polished to 1 mm on both surfaces to test the in-line transmittance. Moreover, to analyze the density trajectories the green bodies were sintered at different temperatures from 1500 1C up to 1750 1C for different time periods from 0.5 h to 20 h. The polished samples were thermally etched in air at 1400 1C or 1500 1C (depending on the sintering temperature) for 3 h to expose the grain boundaries so as to observe the microstructure evolutions. The microscopic morphologies of the raw materials and powder mixtures were observed by a field emission scanning electron microscope (FESEM, Model S-4800,Hitachi, Japan).

2. Experimental High purity commercial powders α-Al2O3 ( 4 99.99%, Alfa Aesar, USA), Y2O3 (4 99.99%, Alfa Aesar, USA), and Nd2O3 (4 99.99%, Alfa Aesar, USA) were used as raw materials. The powders were weighed and mixed in stoichiometric proportions of 1 at% Nd:YAG (Nd0.03Y2.97Al5O12) with 0.08 wt% MgO (4 99.99%, Alfa Aesar, USA) and 0.8 wt% TEOS

Table 1 Summary of particle sizes of the raw materials.

DSEM (nm) D50 (nm) DBET (nm)

Y2O3

α-Al2O3

Nd2O3

3000 1937 484

300 318 157

1000 963 105

Fig. 1. FESEM micrographs of the raw powders (a) Y2O3, (b) α-Al2O3 and (c) Nd2O3.

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Fig. 2. FESEM micrographs of the powder mixtures ball milled for (a) 2 h, (b) 4 h, (c) 8 h, (d) 12 h, (e) 16 h, and (f) 20 h.

Table 2 Specific surface area (SBET) and average particle size (D50) of powder mixtures with different ball milling times. Ball milling time (h) 2

SBET (m /g) D50 (nm)

2

4

8

12

16

20

5.93 508

6.35 485

6.73 497

7.12 459

7.25 459

7.61 400

Fig. 3. Photograph of Nd:YAG ceramics (Φ16 mm  1 mm) sintered at 1760 1C for 3 h from the powder mixtures ball milled for different time periods.

The particle size distributions of the raw materials and the powder mixtures were measured by a Zeta Potential Analyzer (Model Zetaplus, Brookhaven, New York, USA). The specific surface areas of the raw materials and the powder mixtures were measured by a surface area and porosity analyzer (Norcross ASAP 2010, Micromeritics, USA) using the BET multipoint

method. The average particle size (DBET) was calculated from the specific surface area (SBET). Densities of the sintered specimens were measured by the Archimedes method, using deionized water as the immersion medium. Thermally etched surfaces of the mirror-polished ceramics were observed by using a field emission scanning electron microscope (Model JSM6700, JEOL,Japan). The grain sizes of the samples were obtained by the linear intercept method [40] (more than 200 grains were counted), and the average grain size was calculated by multiplying the average linear intercept distance by 1.56. The in-line transmittances of the samples were tested by a UV– vis–NIR spectrophotometer (Model Cary-5000, Varian, USA).

3. Results and discussion Fig. 1 shows the FESEM morphologies of the raw materials. The average particle sizes (DSEM) of Y2O3, α-Al2O3 and Nd2O3 powders are about 3 μm, 0.3 μm and 1 μm, respectively. The particle sizes of the raw powders measured by different methods are summarized in Table 1. The average diameter (D50) was obtained from the particle size distribution, and the equivalent BET diameter (DBET) was calculated from

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where ρ is the density of the material, and SBET is the specific surface area. The densities of Y2O3, Al2O3 and Nd2O3 are 5.01, 3.97 and 7.24 g/cm3, respectively. And the measured specific surfaces of Y2O3, Al2O3 and Nd2O3 are 2.4743, 9.6459 and 7.896 m2/g, respectively. It is worth mentioning the fact that in all the powders DBET is much smaller than visible size D50 of the particles, which means that the particles are porous and consist of small crystallites. Fig. 2 shows the FESEM micrographs of the powder mixtures ball milled for 2 h, 4 h, 8 h, 12 h, 16 h and 20 h, respectively. It can be seen that the coarse particles can be gradually ground into fine particles with prolonging the ball milling time, which can be confirmed by the results shown in Table 2. Fig. 3 shows the photograph of Nd:YAG ceramics (Φ16 mm  1 mm) sintered at 1760 1C for 3 h from powder mixtures ball milled for 2–20 h. When the ball milling time is less than 10 h, the optical transparency of the prepared Nd:YAG ceramic increases gradually with the increase of milling time, which can be clearly observed with the naked eyes. For the samples with ball milling time longer than 10 h, the transparency can be distinguished by a spectrophotometer.

Fig. 4 shows the in-line transmittances of the Nd:YAG ceramics sintered at 1760 1C for 3 h from powder mixtures ball milled for different time periods. The transmittance increases with the increase of ball milling time before reaching the highest transmittances of 83% at 1064 nm and 82% at 400 nm, when the ball milling time is 12 h. And the transmittance shows a slight decrease with further increase of the ball milling time, though the values still remain at the high level. This may be caused by either impurities or supplementary amount of alumina from jar and balls abraded during the milling process. It should be mentioned that 1760 1C–3 h-sintered samples obtained from powder mixtures ball milled for more than 10 h have relative density higher than 99.9%. To investigate the details of the sintering process, relative density trajectories and microstructure evolution of Nd:YAG ceramics at different sintering temperatures and time were studied. Fig. 5 shows the relative density during intermediate and final densification stage as a function of sintering temperature and time. It is noticed that the relative density of each sample increases with the increase of sintering temperature and holding time. For all the samples sintered below 1600 1C, regardless of the holding time, the density increases with increasing ball milling time, as shown in Fig. 5(a). However, when the temperature is higher than 1600 1C, the densification rate shows a slight decrease with ball milling time of more than 12 h. It can be seen from Fig. 5(b) that the sample made from

Fig. 4. In-line transmittances of Nd:YAG ceramics sintered at 1760 1C for 3 h from powder mixtures ball milled for different time periods: (a) 190–1100 nm, (b) 1064 and 400 nm.

Fig. 5. (a) Relative density versus sintering time at 1550 1C for Nd:YAG ceramics from powder mixtures ball milled for different time periods. (b) Relative density versus sintering temperature for Nd:YAG ceramics from powder mixtures ball milled different time periods; the holding time was 2 h.

the following formula: DBET ¼

6 ρSBET

ð1Þ

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Fig. 6. FESEM micrographs of Nd:YAG ceramics sintered at 1600 1C for 2 h from powder mixtures ball milled for (a) 2 h, (b) 4 h, (c) 6 h, (d) 8 h, (e) 10 h and (f) 12 h.

The ball milling time, which determines the particle size, is closely related to the final density due to the sensitivity of different diffusion processes to particle size. As described in Ref. [41], neck growth between contacting particles is very common in sintering process and equations of the neck growth for each transport mechanism usually depend on the particle size. Surface and grain boundary diffusions are most sensitive to particle size, the second is evaporation–condensation, and the least is volume diffusion. Particles with smaller size have higher interface content per unit volume which favors the interfacial diffusion processes. According to the initial mass diffusion kinetics equation 3 1

Fig. 7. Grain size for samples sintered at 1600 1C for 2 h from powder mixtures ball milled for different time periods.

the powder mixture ball milled for 12 h presents the highest density at 1750 1C, which is consistent with the in-line transmittance shown in Fig. 4.

x=r ¼ Kr 5 t 5

ð2Þ

where x/r is the neck size ratio, r is the particle size, t is the time, and K is a constant when the temperature is fixed. The neck size ratio is in direct proportion to  3/5 of the power of the particle size r. This means that powders with larger size can hardly achieve high density even with long sintering time,

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while smaller particles will demonstrate a rapid densification. And the equation explains why the ceramics sintered at the same temperature and time tend to have higher relative density at the initial and intermediate stages of sintering if the green body is made from the powders milled for longer time, as shown in Fig. 5. The relation of relative density with sintering temperature and time at the intermediate and final stages of sintering can be explained using the equation proposed by Coble [42]. At the intermediate and final stages, the relation between porosity Pc and sintering time t is Pc ¼

ADn Ωγ ðt f  tÞ KTL3

ð3Þ

where Dn is the self-diffusion coefficient, K is the sintering rate constant, Ω is the pore volume, γ is the interfacial energy, tf is the complete sintering time, L is the pore length, and A is a constant which is different at the intermediate and final stages of sintering. It is pointed that the porosity decreases with the

increase of sintering time; in other words, the density increases with the sintering time. The sintering temperature has a more complicated effect than the sintering time on the density. The self-diffusion coefficient Dn and the sintering rate K has a relationship with the temperature T as follows: Dn p exp ð

 Ej Þ kB T

K p exp ð Q=RTÞ

ð4Þ ð5Þ

where Ej is the activation energy to jump to an adjacent vacancy, kB is Boltzmann's constant, Q is the sintering activation energy, R is the molar gas constant. Combining Eqs. (3)–(5), it can be highlighted that the porosity decreases with the increase of temperature, which means that the relative density increases with the increase of temperature. Microstructures of Nd:YAG ceramics sintered at 1600 1C for 2 h from powder mixtures ball milled for different time

Fig. 8. FESEM micrographs of ceramics sintered at 1760 1C for 3 h from powder mixtures ball milled for (a) 2 h, (b) 4 h, (c) 8 h, (d) 12 h, (e) 16 h and (f) 20 h.

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periods are shown in Fig. 6. It can be noticed that prolonged ball milling time leads to the decrease of porosity, which is consistent with the result shown in Fig. 5. The grain size increases from 1.06 μm to 1.50 μm, 1.62 μm, 1.76 μm, 2.02 μm, and 2.23 μm with increase of ball milling time, as shown in Fig. 7. Fig. 8 presents FESEM micrographs of Nd:YAG ceramics sintered at 1760 1C for 3 h from powder mixtures ball milled for 2 h, 4 h, 8 h, 12 h, 16 h and 20 h. It can be seen that the number of residual pores gets fewer with increasing ball milling time up to 12 h. The samples obtained from particles with larger size can hardly achieve full density even when sintered at high temperature. The porosity is kept almost unchanged with further increase of ball milling time except that very few residual pores were observed in the sample from powder mixture ball milled for 20 h, as shown in Fig. 8(f). There is no significant grain growth with different ball milling times. The average grain sizes are 13.39 μm,

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13.73 μm, 13.81 μm, 13.42 μm, 13.54 μm and 14.07 μm for 4 h, 8 h, 12 h, 16 h and 20 h, respectively. Although no secondary phase was observed in all the samples, slight deviation from the stoichiometric composition and small amount of residual micropores in Nd:YAG ceramics may influence the optical transmittance, as confirmed in Fig. 4. It must be mentioned that the scratches on the surfaces were caused by the imperfect polishing, as shown in Fig. 8(c) and (f). Fig. 9 presents FESEM micrographs of ceramics sintered at 1550 1C for 0.5–20 h from powder mixtures ball milled for 12 h. For the sample sintered for 0.5 h, lots of micro-pores gathered at grain boundaries forming a pore much larger than the grain size. With the increase of holding time, the pores were gradually removed and separated from the densified regions. However, the closed pores along with few open pores still can be observed even the holding time is as long as 20 h. The grain size increased slightly from 0.58 μm to 0.77 μm,

Fig. 9. FESEM micrographs of ceramics obtained from powder mixtures ball milled for 12 h, sintered at 1550 1C for (a) 0.5 h, (b) 1 h, (c) 3 h, (d) 5 h, (e) 10 h and (f) 20 h.

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Fig. 11(c)), and the last is full dense bodies with an average grain size larger than 2 μm (Fig. 11(d)–(f)). In the last case, a few intragranular pores were observed. These observations show that the final stage of the ceramic densification (ρ 4 96%) is accompanied by significant grain growth. Comparing Figs. 9 and 11, it can be highlighted that the microstructure evolution is more sensitive to temperature than holding time. Full density can be hardly achieved even after sintering for long time (20 h) at low temperature (1550 1C), which coincides with the results in Fig. 5. Fig. 12 presents the sintering trajectory of the samples sintered at 1500–1750 1C for 0.5–20 h from the powder mixtures ball milled for 12 h. It can be seen that the grain size increases slowly (from 0.58 μm to 2.23 μm) until the relative density reaches 97% and becomes very sensitive at higher relative densities, which coincides with the results shown in Figs. 9–11. To determine the grain growth exponent (n) and rate constant (k), the kinetic value can be calculated from the following equation: Gn  Gn0 ¼ kt

Fig. 10. Grain sizes for samples obtained from powder mixtures ball milled time for 12 h (a) sintered at 1550 1C for different time periods and (b) sintered at different sintering temperatures for 2 h.

1.16 μm, 1.36 μm, 1.97 μm, and 2.22 μm for the holding time from 0.5–1 h, 3 h, 5 h, 10 h, and 20 h, respectively, as shown in Fig. 10(a). Fig. 11 presents FESEM micrographs for ceramics sintered at different temperatures for 2 h from powder mixtures ball milled for 12 h. The sample sintered at 1500 1C shows quite a lot of pores gathered forming larger ones than the grains which are similar to the sample described above in Fig. 9(a). The pores were removed rapidly with the increase of temperature, and no significant pores remained at the temperature of 1700 1C and higher. Grain size increased rapidly from 0.70 μm to 1.11 μm, 2.23 μm, 6.08 μm, 6.51μ m, and 8.54 μm when the temperature increased from 1500 1C to 1550 1C, 1600 1C, 1650 1C, 1700 1C, and 1750 1C, as shown in Fig. 10(b). Combination of Figs. 9 and 11 presents three types of microstructures as a function of the two sintering parameters (sintering temperature and time) which is similar to the results of Boulesteix et al. [43]. One is porous ceramics with open porosities and the relative density ranging from 80% to 90% with a micrometer grain size under or around 1 μm (Fig. 9(a)–(c) and Fig.11(a) and(b)); other is dense ceramics with closed intergranular porosity, relative density ranging from 90% to 99% with a micrometer grain size under or around 2 μm (Fig. 9(d)–(f) and

ð6Þ

where G is the average grain size, G0 is taken to be the grain size at 0.5 h at the temperature of 1550 1C, and t is the holding time. The grain growth exponent n ranges from 2–4, and 3 fits the best for the samples in this paper, which is consistent with Stevenson's results [31]. And 1550 1C is just in the temperature range of 1525–1850 1C which was proved by Boulesteix et al. [44] and Kochawattana et al. [32] that the grain growth kinetics fit a cubic rate law. The grain growth kinetics curve with the R2 ranging from 0.97–0.99 is shown in Fig. 13. The k value calculated from the curve is about 1.53E  22 m3/s. It should be mentioned that the value at 10 h was probably caused by the measurement deviation.

4. Conclusions Nd:YAG ceramics were prepared by solid-state reactive sintering of oxides powder mixtures ball milled for different hours. The coarse powders can be ground to smaller particles with increase of ball milling time. Increasing ball milling time benefits the densification process because the diffusion process is very sensitive to the particle size. Smaller particle size leads to shorter sintering time or lower temperature compared with larger ones to achieve an equivalent degree of sintering. Porosity decreases by prolonging the ball milling time up to 12 h and is kept almost unchanged with further increase of time. Slight deviation from the stoichiometric composition and small amount of residual micro-pores in Nd:YAG ceramics caused by the overlong ball milling time will decrease the optical transmittance of the ceramics. With the optimal ball milling process, Nd:YAG transparent ceramics with the in-line transmittance of 83% at 1064 nm were obtained by sintering at 1760 1C for 3 h from the powder mixture ball milled for 12 h.

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Fig. 11. FESEM micrographs of ceramics obtained from powder mixtures ball milled for 12 h, sintered for 2 h at (a) 1500 1C, (b) 1550 1C, (c) 1600 1C, (d) 1650 1C, (e) 1700 1C and (f) 1750 1C.

Fig. 12. Sintering trajectory for Nd:YAG ceramics sintered between 1500 1C and 1750 1C for 0.5–20 h from powder mixtures ball milled for 12 h.

Fig. 13. G3 G30 vs sintering time at 1550 1C for ceramics obtained with powder mixture ball milled for 12 h.

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Acknowledgments This work was supported by the Project of International Cooperation and Exchange NSFC-RFBR (Grant nos. 51311120078 and 512111056/13-02-91173-ГФЕН_а), the Key Program of National Natural Science Foundation of China (Grant no. 91022035) and the Project for Young Scientists Fund of National Natural Science Foundation of China (Grant nos. 51002172 and 51302298). References [1] A. Ikesue, Y.L. Aung, Ceramic laser materials, Nat. Photonics 2 (2008) 721–727. [2] A. Ikesue, Y.L. Aung, T. Taira, T. Kamimura, K. Yoshida, G.L. Messing, Progress in ceramic lasers, Annu. Rev. Mater. Res. 36 (2006) 397–429. [3] J. Li, Y.B. Pan, Y.P. Zeng, W.B. Liu, B.X. Jiang, J.K. Guo, The history, development, and future prospects for laser ceramics: a review, Int. J. Refract. Met. Hard Mater. 39 (2013) 44–52. [4] A. Ikesue, Polycrystalline Nd:YAG ceramics lasers, Opt. Mater. 19 (2002) 183–187. [5] A. Ikesue, T. Kinoshita, K. Kamata, K. Yoshida, Fabrication and optical properties of high-performance polycrystalline Nd:YAG ceramics for solid-state lasers, J. Am. Ceram. Soc. 78 (1995) 1033–1040. [6] J.R. Lu, M. Prabhu, J.Q. Xu, K. Ueda, H. Yagi, T. Yanagitani, A.A. Kaminskii, Highly efficient 2% Nd:yttrium aluminum garnet ceramic laser, Appl. Phys. Lett., 77, , 2000, p. 3707–3709. [7] J.R. Lu, K. Ueda, H. Yagi, T. Yanagitani, Y. Akiyama, A.A. Kaminskii, Neodymium doped yttrium aluminum garnet (Y3Al5O12) nanocrystalline ceramics—a new generation of solid state laser and optical materials, J. Alloys Compd. 341 (2002) 220–225. [8] J. Li, Y.S. Wu, Y.B. Pan, W.B. Liu, L.P. Huang, J.K. Guo, Fabrication, microstructure and properties of highly transparent Nd:YAG laser ceramics, Opt. Mater. 31 (2008) 6–17. [9] S.H. Lee, S. Kochawattana, G.L. Messing, Solid-state reactive sintering of transparent polycrystalline Nd:YAG ceramics, J. Am. Ceram. Soc. 89 (2006) 1945–1950. [10] W.B. Liu, J. Li, B.X. Jiang, D. Zhang, Y.B. Pan, 2.44 KW laser output of Nd:YAG ceramic slab fabricated by a solid-state reactive sintering, J. Alloys Compd. 538 (2012) 258–261. [11] L. Esposito, A. Piancastelli, Role of powder properties and shaping techniques on the formation of pore-free YAG materials, J. Eur. Ceram. Soc. 29 (2009) 317–322. [12] Y.S. Wu, J. Li, Y.B. Pan, Q. Liu, J.K. Guo, B.X. Jiang, J. Xu, Diodepumped passively Q-switched Nd:YAG ceramic laser with a Cr4 þ :YAG crystal-satiable absorber, J. Am. Ceram. Soc. 90 (2007) 1629–1631. [13] J.G. Li, T. Ikegami, J.H. Lee, T. Mori, Y. Yajima, Co-precipitation synthesis and sintering of yttrium aluminum garnet (YAG) powders: the effect of precipitant, J. Eur. Ceram. Soc. 20 (2000) 2395–2405. [14] H.Z. Wang, L. Gao, K. Niihara, Synthesis of nanoscaled yttrium aluminum garnet by the co-precipitation method, Mater. Sci. Eng. A 288 (2000) 1–4. [15] X. Li, Q. Li, J.Y. Wang, S.L. Yang, H. Liu, Synthesis of Nd3 þ doped nano-crystalline yttrium aluminum garnet (YAG) powders leading to transparent ceramic, Opt. Mater. 29 (2007) 528–531. [16] W. Zhang, T.C. Lu., B.Y. Ma, N. Wei, Z.W. Lu, F. Li, Y.B. Guan, X.T. Chen, W. Liu, L. Qi, Improvement of optical properties of Nd:YAG transparent ceramics by post-annealing and post hot isostatic pressing, Opt. Mater. 35 (2013) 2405–2410. [17] R.P. Yavetskiy, E.A. Vovk, A.G. Doroshenko, M.I. Danylenko, A.V. Lopin, I.A. Petrusha, V.F. Tkachenko, A.V. Tolmachev, V.Z. Turkevich, Y3Al5O12 translucent nanostructured ceramics-obtaining and optical properties, Ceram. Int. 37 (2011) 2477–2484. [18] I.S. Puzyrev, I.V. V’yukhina, M.G. Ivanov, Yu.G. Yatluk, Development of methods for preparation of Nd:YAG nanoparticles, Glass Phys. Chem. 38 (2012) 427–430.

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