Co-precipitation synthesis and vacuum sintering of Nd:YAG powders for transparent ceramics

Materials Research Bulletin 70 (2015) 365–372

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Co-precipitation synthesis and vacuum sintering of Nd:YAG powders for transparent ceramics W. Zhang a,b, * , T.C. Lu a,c, N. Wei a,c , Y.L. Shi a , B.Y. Ma d , H. Luo a,b , Z.B. Zhang b , J. Deng b , Z.G. Guan b , H.R. Zhang b , C.N. Li b , R.H. Niu b a

College of Physical Science and Technology, Sichuan University, 610064 Chengdu, China Southwest Institute of Technical Physics, Chengdu 610041, China Key Laboratory of High Energy Density Physics and Technology of Ministry of Education, Sichuan University, Chengdu 610064, China d College of Physics and Electronic Engineering, Nanyang Normal University, Nanyang 473061, China b c

A R T I C L E I N F O

A B S T R A C T

Article history: Received 10 November 2014 Received in revised form 22 April 2015 Accepted 30 April 2015 Available online 5 May 2015

The Nd:YAG nanopowders were synthesized at 800–1300
C via a modified co-precipitation method. Nd: YAG transparent ceramics were fabricated by vacuum sintering at 1770–1790
C for 20 h with heating rates 1–10
C/min. The effects of calcining temperatures of highly sinterable Nd:YAG nanopowders, sintering temperature and heating rate on microstructure and transmittance of Nd:YAG ceramics have been investigated. After calcination at 1250
C for 4 h, Nd:YAG nanopowders with primary particle size of about 106 nm and low agglomeration were obtained. Using as prepared powders, high optical quality Nd: YAG ceramics prepared at 1780
C with dwell time of 20 h and heating rate at 1
C/min have been obtained and shows the in-line transmittance reached 83.9% (1064 nm). With 17.6 W of maximum absorbed pump power, laser output of 4.5 W has been obtained with an oscillation threshold and a slope efficiency of 0.5 W and 25.6%. ã 2015 Elsevier Ltd. All rights reserved.

Keywords: Nd:YAG powders Transparent ceramics Co-precipitation synthesis Optical properties Sintering

1. Introduction During the last 30 years, neodymium-doped yttrium aluminum garnet (i.e., Y33xNd3xAl5O12, Nd:YAG) crystal has been proven to be a very good solid-state laser material, in particular for power applications thanks to its good optical and thermomechanical properties. Low duration and low cost of manufacturing, absence of severe limitation in size and geometry make the Nd:YAG ceramics attractive with comparison to single crystals [1–3]. Recently, polycrystalline Nd:YAG ceramic laser materials have received much attention because the optical quality has been improved greatly and highly efficient laser oscillations could be obtained that comparable in efficiency with Nd:YAG crystals [4–10]. But it was reported that the optical quality of the conventional transparent ceramics was a little inferior to that of single crystal material, because of the residual pores and grain boundaries in the ceramics acting as optical scattering centers. In order to decrease optical loss, it is extremely important to fabricate highly dispersed ultra-fine YAG powders and considerably full

* Corresponding authors at: Southwest Institute of Technical Physics, Chengdu 610041, China. Tel: +86 2885222120; fax: +86 2885222120. E-mail address: [email protected] (W. Zhang). http://dx.doi.org/10.1016/j.materresbull.2015.04.063 0025-5408/ ã 2015 Elsevier Ltd. All rights reserved.

dense ceramics with pore-free structure and clean grain boundaries. The history of the research on YAG transparent ceramics dates back to the end of last century, Ikesue et al. fabricated the highly quality YAG ceramics by solid-state reaction method first [4]. Solidstate reactive sintering is a relatively simple way to fabricate YAG transparent ceramics and a range of compositions is easy to implement by changing the reactant powder amounts during batching. However, the main disadvantage of this process is that the incorporation of some impurities is unavoidable during ball milling. Another disadvantage is that it requires long time and very high temperature. Fortunately, wet chemical approaches such as co-precipitation [11,12], spray pyrolysis [13], hydrothermal synthesis [14] and sol–gel combustion synthesis [15] have many advantages such as atomic level mixing of high-purity precursors and low processing temperature. Among them, co-precipitation is one of the most promising techniques for the preparation of YAG nanopowders with excellent chemical homogeneity, good crystallinity and pure phase at low temperature [11]. But the synthesis process and the sinterability of powders still need to be improved. The purpose is to fabricate fully dense and transparent ceramics by vacuum sintering of YAG nanopowders with high sinterability. On the other hand, understanding the links between sintering parameters and microstructure evolution is important to master

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screen. The sieved precursors were then calcined in air at temperatures of 800, 900, 1000, 1200, 1250 and 1300
C for 4 h at a heating rate of 10
C/min to form garnet phase of YAG powders. The phase development of the calcined powders was observed by X-ray diffractometer (XRD, Cu-Ka radiation 1.54 Å, D/max-rA, Rigaku, Japan). The average crystallite size of the calcined powders was calculated from X-ray peak broadening using Scherrer formula. The XRD data are refined by Rietveld method using the Jade program (Version 6.5). The morphology and dispersion of the powders were observed using scanning electron microscopy (SEM, Model S-4800, Hitachi, Japan). The average size over the number of particles were manually estimated by measuring diameters grains L in the SEM photographs of 400 particles, and the average size D, was determined by D = 1.558L [30].

optical properties of transparent Nd:YAG ceramics. Generally, in the Nd:YAG ceramics sintering procedure, the average grain size, residual pores, and other defects of the ceramics mainly depend on the parameters of sintering [16,17]. As a result, study on sintering parameters of YAG transparent ceramics has become a hot research topic [18], in terms of reducing of pores and other defects. There are many reports regarding sintering parameters and related mechanism of Nd:YAG ceramics by solid-state sintering of oxide powders [16–18]. In these cases, sintering temperature, holding time, heating rate, sintering aids and rare-earth diffusion were investigated [19–24]. Meanwhile, Li et al. have reported that the sintering temperature of Nd:YAG ceramics with co-precipitation nanopowders were sintered into transparent bodies by vacuum sintering at 1730–1790
C [10]. Nevertheless, the sintering temperature and heating rates as well as underlying mechanism of Nd:YAG ceramics using co-precipitation nanopowders, was not investigated systematically so far. We aim to develop an appropriate sintering processing conditions to produce high transparent Nd:YAG ceramics with co-precipitation nanopowders of high sinterability. In this work, we synthesized 2 at.% Nd:YAG nanopowders at 800–1300
C through a modified co-precipitation method. Subsequently, using the prepared powders calcined at 1250
C, Nd:YAG transparent ceramics were fabricated by vacuum sintering at 1770–1790
C for 20 h with heating rates 1–10
C/min and then annealed at 1450
C for 20 h in air. We investigated how calcining temperature affects sinterability of Nd:YAG nanopowders as well as heating rate and sintering temperature influence the optical properties and microstructure of the Nd:YAG transparent ceramics. Furthermore, we discussed the detailed mechanism of related sintering treatments.

2.2. Ceramic fabrication and characterization Using the powder calcined at 1250
C as starting materials, the Nd:YAG nanopowders were blended with alumina balls (99.99% Al2O3, diameter approx. 1 cm) for 24 h in ethanol using 0.5 wt.% tetraethyl orthosilicate (TEOS) as a sintering aid. Then, the alcohol solvent was removed by drying the milled slurry in oven, and residual organic material was removed by calcining at 800
C for 20 h in flowing oxygen. The Nd:YAG nanopowders were uniaxially pressed into Ø 20–50 mm disks at 20 MPa and then isostatically cold pressed at 200 MPa. Then the compacted disks were sintered at 1770–1790
C for 20 h in a graphite-heated vacuum furnace under vacuum (1 103 Pa) with 1–10
C/min heating rates. Thus, the samples were obtained and hereafter named as S1, S2, S3, S4, S5, S6 and S7, respectively, as shown in Table 1. After cooling, the sintered samples were air annealed at 1450
C for 20 h to eliminate oxygen vacancies formed during vacuum sintering. The phase structures of sintered ceramics were identified by X-ray diffraction (XRD, Model D/max-rA, Rigaku, Japan). The ceramic samples were ground to the thickness of 1.2 mm or 2 mm and thoroughly polished carefully on both sides to eliminate surface scattering. The real in-line transmittances of samples were measured using an ultraviolet spectroscopy spectrometer (Spectrophotometer, Model UV-1700 Pharmaspec, Shimadzu, Japan) in the wavelength from 200 and 1100 nm. Before microstructural observations, sintered samples were previously thermally etched under air during 2 h at temperatures lower than sintering one by 300
C. Grain size was measured thanks to image analysis performed on micrographs obtained using SEM (Model S-4800, Hitachi, Japan). For each sample, average grain size measurements carried out over 400 grains via the lineal analysis intercept technique using the factor 1.558 times the average intercept length. In order to determine pores using the SEM quantitatively, we chose 25 images (each measurement image 100  100 mm2) that were located in an arrangement of a 5  5 lattice with an interval of 300 mm between each other in an area with a scale of 1700  1700 mm2 of every Nd:YAG sample surface. Subsequently

2. Experimental 2.1. Powder synthesis and characterization The manufacturing process during this work is similar to that reported in our previous works [25–29]. Neodymium nitrate (Nd (NO3)36H2O, 99.99%), yttrium nitrate (Y(NO3)36H2O, 99.99%) and ammonium aluminum sulfate (NH4Al(SO4)212H2O, 99.99%) were mixed together in stoichiometric proportions to form Nd0.06Y2.94Al5O12 (2 at.% Nd:YAG) in distilled water. The precipitant solution was prepared by dissolving ammonium hydrogen carbonate (NH4HCO3, analytical grade) in mixed solvent of alcohol and distilled water. The volume ratio of alcohol to distilled water is 0.6. The mixed solution was dripped into the precipitant solution at a dripping speed of 3 ml/min under stirring at 18
C. With the suspension aged for 24 h, filtered and washed repeatedly with distilled water and alcohol in sequence, the precipitate was obtained. Then precursors were produced after the precipitate was dried at 80
C for 24 h with an oven. The dried cake was crushed with an agate pestle and mortar, and sieved through a 200-mesh

Table 1 The sizes, sintering conditions, average grain size, average pore size and porosity of the sintered ceramic samples. Samples

Size

Sintering condition

S1 S2 S3 S4 S5 S6 S7

Grain size

Pore size

Porosity

(mm )

Temperature ( C)

Heating rate ( C/min)

Dwell time (h)

Gg (mm)

Dp (mm)

Vp (%)

Ø16.8  1.2 Ø16.8  1.2 Ø16.8  1.2 Ø16.8  2 Ø16.8  2 Ø16.8  2 Ø40  2

1770 1780 1790 1780 1780 1780 1780

10 10 10 10 2 1 1

20 20 20 20 20 20 20

9.99 10.17 10.21 9.03 9.12 9.26 8.85

0.28 – 0.31 0.52 0.36 – –

0.0023 – 0.0016 0.0027 0.0014 – –

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smaller images (10  10 mm2) were taken over the pores one by one until all the pores in the larger scan were analyzed in every image. The sizes of each pore in all images were manually defined by the MB-ruler software so that the area on the scale section could be counted. 2.3. Laser experiment The schematic diagram of the laser experimental setup is shown in Fig. 1. The laser cavity was an end-pumped plano–plano resonator with total cavity length of 35–40 mm. Fiber coupled 808 nm diode-lasers with core diameter of 400 mm were used as the pump sources with C.W. laser operation. Two convex lenses were used to focus the pump beam into the ceramic sample and to produce the pump light footprint of about 300 mm. The total cavity was mainly composed of one input mirror and one output coupler. The input mirror had 95% transmission at pumping wavelength and 99.8% reflectivity at laser wavelength. The output coupler mirrors with transmission of 10% for Nd:YAG ceramics were used to measure 1064 nm laser performance. The ceramic was placed as close as possible to the input mirror and held on the brass mount without special heat treatment. 3. Results and discussion 3.1. X-ray diffraction and microstructural of Nd:YAG powders Fig. 2 shows the XRD patterns of as-synthesized precursor and the powders calcined at different temperatures from 800
C to 1300
C for 4 h. The Nd3+ ions enter into the crystal and occupy the Y3+ lattice sites. Because the radius of Nd3+ ions is a little larger than that of Y3+ ions (r (Nd3+) = 1.12 Å, r (Y3+) = 1.02 Å), which can cause the lattice distortion and the lattice parameter of Nd:YAG samples slightly larger than the standard value of cubic YAG (12.0089 Å) [31]. Since no obvious diffraction peaks are observed, it can be concluded that the precursor was amorphous up to 800
C. The crystalline YAG phase can be obtained at a higher temperature (900
C). A small amount of the intermediate phase, YAlO3 (YAP), was detected at 900
C. The result is different from the conclusion reached in previous work [32,33]. The discrepancy is likely to be caused by inhomogeneity of Y3+ and Al3+ ions in the precursors in this experiment. YAG is the only phase detected at and above 1000
C, and more distinct peak shapes with stronger intensities were observed, indicating crystallite growth in the YAG powder with increasing temperature. The average crystallite size of YAG powder defined as was determined by X-ray line broadening and calculated using the Scherrer equation: DXRD ¼

0:89l Bcosu

(1)

where B = (BO  BC) ,BO is the full width at half maximum (in
2u), BC is the correction factor for instrument broadening, u is the angle of the peak maximum (in
2u ), and l is the CuKa weighted average wavelength. For the powders calcined at 1000, 1100, 1200, 2

Fig. 2. XRD patterns of as-synthesized precursor and the resultant powders obtained by calcining the precursors at different temperatures for 4 h.

1250 and 1300
C, the mean crystallite size values of YAG powders calculated from the (4 2 0) planes are 31.7 nm, 35.3 nm, 37.5 nm, 105.6 nm, 112.3 nm, respectively, as listed in Table 2. Fig. 3 shows the SEM micrographs of the powders from the precursor calcined at different temperatures for 4 h. The morphology of the as-synthesized precursor is depicted in Fig. 3(a). The precursor mainly contains sub-micrometer sized aggregates of nano-sized primary particles. The resultant powders calcined at 800 and 900
C were agglomerated and showed similar overall morphology to that of the precursor. For the powders calcined at 1000, 1100, 1200, 1250 and 1300
C, appreciable particle growth from 108 to 382 nm occurred by increasing the calcination temperature. The mean primary particle size estimated from SEM photographs is denoted as DSEM. It can be found that the DSEM value (175 nm) of the Nd:YAG powder calcined at 1250
C is just a little larger than that of DXRD (105.6 nm). This indicates that each particle shown in the SEM photograph is almost a single crystallite. Moreover, the primary particles with mean diameter of 175 nm are weakly agglomerated by small connections with necks. According to the above analysis, it is considered that the resultant powder calcined at 1250
C will be favorable for transparent ceramic fabrication. Higher calcination temperature caused drastic increase in crystallite size and decrease in sinterability. While the YAG powders produced at lower temperature have smaller crystallite sizes, they had more secondary phases and exhibited lower densification rates [12]. The Nd:YAG powder calcined at various temperatures (800–1300
C) was sintered to transparent ceramics which would be reported in our future paper.

Table 2 The calcining conditions, average crystalline sizes and particle sizes of the experimental precursor and powders. Calcining conditions

Fig. 1. Schematic diagram of the end-pumping laser experimental setup. (L1, L2: focusing lens; M1: cavity mirror; M2: output coupler; M: the total reflectivity at laser wavelength).

Average crystalline size

Particle size

Temperature (
C)

Dwell time (h)

DXRD (nm)

DSEM (nm)

Precursor 800 900 1000 1100 1200 1250 1300

– 4 4 4 4 4 4 4

– – – 31.7 35.3 67.5 105.6 112.3

8.7 8.9 8.9 108 115 169 175 382

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Fig. 3. SEM micrographs of the starting powders calcined at different temperatures for 4 h: (a) precursor; (b) 800
C; (c) 900
C; (d) 1000
C; (e) 1100
C; (f) 1200
C; (g) 1250
C and (h) 1300
C.

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Fig. 4. Photograph of Nd:YAG bodies sintered at (S1) 1770
C, (S2) 1780
C and (S3) 1790
C (a) and the in-line transmittances of the samples (b).

3.2. Effect of sintering temperature The in-line transmittance is the main parameter for evaluating the optical properties of Nd:YAG transparent ceramics. The optical in-line transmittances and photograph of the 1.2 mm-thick transparent Nd:YAG ceramics sintered under vacuum for 20 h at different temperatures in the range 1770–1790
C are shown in Fig. 4. It should be noteworthy that the samples S1, S2 and S3 have high transparency and the letters under the ceramics can be seen distinctly. Interestingly, the UV absorption edges of samples shift to 270 nm. As we reported previously, the red shift after vacuum sintering and air annealing is attributed to the absorption by Fe3+ charge transfer bands, and iron impurities were from raw materials and preparation process [25]. This process could be expressed as follows:

FenFe

Reducingatmosphere 1  1 ! þ OO $ Fen1 Fe þ VO þ O2 ðgÞ 2 2

passes through transparent ceramics. However, at the wavelength of 300 nm, the optical transmittances of the samples S1, S2 and S3 were 21.3%, 32.1% and 57.2% respectively, depending on the final temperature during vacuum sintering, probably related to the presence of scattering centers (pores and secondary phases) in samples. In order to detect the main reason for the low transmittance of the samples which are lower than the theoretical value, the phase compositions of the samples are identified by XRD, as shown in

(2)

Oxidizingatmosphere

As can be seen in Fig. 4, the transmittance curves of all samples have similar slopes and their transmittances increase with increasing wavelength from 300 to 1100 nm, which originates from existence of scattering centers in the samples. When the optical scattering center is significantly smaller than the wavelengths, the scattering intensity is inversely proportional to the fourth power of the wavelength according to the Rayleigh’s equation [34,35]. The transmittances of the three samples are all around 81% at lasing wavelength of 1064 nm, because the long wavelength is slightly affected by scattering centers when light

Fig. 5. XRD patterns of Nd:YAG bodies sintered at (S1) 1770
C, (S2) 1780
C and (S3) 1790
C.

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Fig. 6. SEM images of thermal-etched surfaces of Nd:YAG bodies sintered at (S1) 1770
C, (S2) 1780
C and (S3) 1790
C. White boxes refer to pores.

Fig. 5. It can be seen that the phase structures of all the samples are almost the same despite of different sintering temperatures. All the diffraction peaks of samples can be well indexed as the cubic garnet structure of Y3Al5O12 (YAG, JCPDS 33-0040), and no other phases or impurities are detected. Moreover, the value of transmission is much lower than the theoretical value due to the existence of residual pores, as can be confirmed in Fig. 6. The microstructures of the thermal etched surfaces of the samples sintered at different temperatures are shown in Fig. 6. From the computerised image analysis and measurements on series of micrographs of this type, the porosity contents (Vp), the size pore (Dp) and the grain size (Gg) are summarised in Table 1. The average grain sizes of the samples ranging from 9.99 mm to 10.21 mm, together with small statistical errors of the measurements themselves, we conclude that there is an insignificant increase in the grain size during vacuum sintering at 1770–1790
C. These microstructures confirm our understanding of residual porosity leading to the difference in in-line transmittance of Nd: YAG samples. In Fig. 6(a), the porosity was not completely removed after vacuum sintering at 1770
C for 20 h. The calculated porosity is order of ppm from the counting method using SEM pictures. Notice that the residual pores, which may be because the rate of grain-boundary was much slower than that of pores mobility during the final stage of lower temperature sintering, was mainly at the triple points of the grains. On the contrary, at the final stage of higher temperature sintering, the grain-boundary migration rate becomes superior to the pore one, the pore and grainboundary plane separate to each other and becomes intragranular [17], as shown in Fig. 6(c). In Fig. 6(b), it can be clearly seen that the nearly pore-free microstructure of the optically transparent ceramics at 1780
C. It indicated that a sintering temperature of 1780
C was adequate to obtain Nd:YAG transparent with high density and a higher transparency. As a consequence, an appropriate sintering temperature should be helpful for improvement of the optical properties of Nd:YAG ceramics.

3.3. Effect of heating rate The 2 mm-thick ceramic samples sintered with different heating rates are shown in Fig. 7. They are clearly transparent. The in-line transmittance is shown in Fig. 8. The transmittance drastically increased with decreasing heating rates from10
C/min to 1
C/min. Compared with the transmittance of samples S4 (10
C/min) and S5 (2
C/min), the samples S6 and S7 sintered with heating rate at 1
C/min showed much higher transmittance. In contrast to the sample S3 (Ø 16.8 mm  1.2 mm), the transmittance

Fig. 7. Photograph of Nd:YAG bodies sintered with heating rate: 10
C/min (S4), 2
C/min (S5), 1
C/min (S6) and 1
C/min (S7).

W. Zhang et al. / Materials Research Bulletin 70 (2015) 365–372

Fig. 8. In-line transmittances of Nd:YAG bodies sintered with heating rate: 10
C/ min (S4), 2
C/min (S5), 1
C/min (S6) and 1
C/min (S7)

of samples S4 (Ø 16.8 mm  2 mm) is lower. This is not surprising, since the different thickness has a strong effect on transmittance of ceramics [36]. Fig. 9 shows microstructures of the sintered samples with various heating rates. The porosity contents (Vp), the pore size (Dp) and the grain size (Gg) of the samples S4–S7 are listed in Table 1. The grain size was 9.03, 9.12, 9.26 and 8.85 mm, for S4 (10
C/min), S5 (2
C/min), S6 (1
C/min) and S7 (1
C/min), respectively. Considering small statistical errors of the measurements and the different size of the samples, we speculate that there is an insignificant change in the grain size during vacuum sintering various heating rates. However, a few distinct pores were detected

371

at the triple points of the grains or within the grains. Especially, submicrometer pores were outstanding for the samples S4 (10
C/min) and S5 (2
C/min), shown in Fig. 9(a) and (b). That should be because increase in the heating rate raises the rapid densification rate of the Nd:YAG ceramics, which obstructs pores exclusion and element diffusion. The results of the present work indicate that these residual pores in Nd:YAG ceramics increase with increasing of heating rate. The optical transmittance of Nd: YAG ceramics is closely related to pores, so the optical properties of samples S6 and S7 are the most excellent among the four samples because of its low porosity. The result also indicates that the lower heating rate is beneficial to improve the optical properties of Nd: YAG ceramics. Interestingly, the pores with the size of several microns present in the samples are likely a result of voids or pressure gradients during green bodies forming process. The best way to eliminate such micro-pores is moving towards a colloidal forming method. Slip casting, gelcasting, tape casting, etc., are all better suited to forming transparent ceramics than dry pressing because they are better to handle nanoscaled and microscaled powders and they produce more homogeneous green bodies [37]. The details of forming techniques of Nd:YAG ceramics will be reported in our future paper. The laser properties of the obtained ceramics were measured. Fig. 10 shows the laser output versus pump power for sample S7 at 1064 nm. With 17.6 W of maximum absorbed pump power, laser output power of samples S7 is 4.5 W and the slope efficiency is 25.6%. The experiment yielded oscillation threshold value of 0.5 W. The laser results show that the quality of the Nd:YAG ceramic sample is inferior to that of single crystal. This may be attribute to the scattering loss mainly caused by a very small quantity of residual pores, which would be too small to be detected by SEM, in the Nd:YAG ceramics. It is believed that higher laser output can be realized by improvement of fabrication details in the further study.

Fig. 9. SEM micrographs of thermal-etched surfaces of samples sintered with heating rate: (a) 10
C/min (S4), (b) 2
C/min (S5), (c) 1
C/min (S6) and (d) 1
C/min (S7). White boxes refer to pores.

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References

Fig. 10. Laser output at 1064 nm as a function of the pump power.

4. Conclusions In this work high transparent Nd:YAG ceramics were successfully obtained by following proper temperature-heating rate sintering schedules. The highly sinterable Nd:YAG nanopowders with primary particle size of about 106 nm and low agglomeration were synthesized at 1250
C for 4 h using a modified co-precipitation method precursor. Using the nanopowders calcinating at 1250
C as prepared powders, sintering parameters were evaluated in relation to microstructure and optical transmission data of sintered samples. Better results show the in-line transmittance of the sample reached 77.8% (400 nm) and 83.9% (1064 nm) respectively, were obtained when prepared at 1780
C with dwell time of 20 h and heating rate at 1
C/min. With 17.6 W of maximum absorbed pump power, laser output of 4.5 W has been obtained with an oscillation threshold and a slope efficiency of 0.5 W and 25.6%. The results indicated that the lower heating rate and appropriate sintering is beneficial to improve the optical properties of Nd:YAG ceramics. The highly sinterable Nd:YAG nanopowders, appropriate sintering temperature and lower heating rate are effective approaches to fabrication of this laser material. Acknowledgements This work was supported by the National Natural Science Foundation of the People’s Republic of China (Grant Nos. 51002098 and 11145006), Science and Technology Project in Shantou (Grant No. 2014SS019), Key Laboratory of Neutron Physics, CAEP (Grant No. 2014BC01) and National High Technology Research and Development Pro-gram (863).

[1] A. Ikesue, Y.A. Aung, Nat. Photonics 2 (12) (2008) 721–727. [2] P. Palmero, B. Bonelli, G. Fantozzi, G. Spina, G. Bonnefont, L. Montanaro, J. Chevalier, Mater. Res. Bull. 48 (2013) 2589–2597. [3] M. Suárez, A. Fernández, J.L. Menéndez, M. Nygren, Z. Zhao, R. Torrecillas, A.-M. Pablo, J.J. Haro-González, I.R. Martín, Ceram. Inter. 40 (2014) 15951–15956. [4] A. Ikesue, T. Kinoshita, K. Kamata, K. Yoshida, J. Am. Ceram. Soc. 78 (1995) 1034–1040. [5] E.H. Penilla, Y. Kodera, J.E. Garay, Mater. Sci. Eng. B 177 (2012) 1178–1187. [6] H. Yagi, T. Yanagitani, K. Takaichi, K.I. Ueda, A.A. Kaminskii, Opt. Mater. 29 (2007) 1258–1262. [7] A. Maître, C. Sallé, R. Boulesteix, J.F. Baumard, R. Boulesteix, Y. Rabinovitch, J. Am. Ceram. Soc. 91 (2) (2008) 406–413. [8] L. Yang, T.C. Lu, H. Xu, W. Zhang, B.Y. Ma, J. Appl. Phys. 107 (2010) 064903. [9] M. Serantoni, A. Piancastelli, A.L. Costa, L. Esposito, Opt. Mater. 34 (2012) 995–1001. [10] J. Li, J. Liu, B.L. Liu, W.B. Liu, Y.P. Zeng, X.W. Ba, T.F. Xie, B.X. Jiang, Q. Liu, Y.B. Pan, X.Q. Feng, J.K. Guo, J. Eur. Ceram. Soc. 34 (10) (2014) 2497–2507. [11] G.G. Xu, X.D. Zhang, W. He, H. Liu, H. Li, R.I. Boughton, Mater. Lett. 60 (2006) 962–965. [12] J. Li, F. Chen, W.B. Liu, W.X. Zhang, L. Wang, X.W. Ba, Y.J. Zhu, Y.B. Pan, J.K. Guo, J. Eur. Ceram. Soc. 32 (2012) 2971–2979. [13] Y.H. Zhou, J. Lin, M. Yu, S.M. Han, S.B. Wang, H.J. Zhang, Mater. Res. Bull. 38 (2003) 1289–1299. [14] Y. Hakuta, T. Haganuma, K. Sue, T. Adschiri, K. Arai, Mater. Res. Bull. 38 (2003) 1257–1265. [15] M. Veith, S. Mathur, A. Kareiva, M. Jilavi, M. Zimmer, V. Huch, J. Mater. Chem. 9 (1999) 3069–3079. [16] R. Boulesteix, A. Maître, J.F. Baumard, C. Sallé, Y. Rabinovitch, Opt. Mater. 31 (2009) 711–715. [17] R. Boulesteix, A. Maître, L. Chrétien, Y. Rabinovitch, C. Sallé, J. Am. Ceram. Soc. 96 (6) (2013) 1724–1731. [18] J. Lu, M. Prabhu, J. Song, C. Li, J. Xu, K. Ueda, A.A. Kaminskii, H. Yagi, T. Yanagitani, Appl. Phys. B 71 (2000) 469–473. [19] W.B. Liu, B.X. Jiang, W.X. Zhang, J. Li, J. Zhou, D. Zhang, Y.B. Pan, J.K. Guo, Ceram. Inter. 36 (2010) 2197–2201. [20] S. Kochawattana, A. Stevenson, S.H. Lee, M. Ramirez, V. Gopalan, J. Dumm, V.K. Castillo, G.J. Quarles, G.L. Messing, J. Eur. Ceram. Soc. 28 (2008) 1527–1534. [21] A.J. Stevenson, X. Li, M.A. Martinez, J.M. Anderson, D.L. Suchy, E.R. Kupp, E.C. Dickey, K.T. Mueller, G.L. Messing, J. Am. Ceram. Soc. 94 (5) (2011) 1380–1387. [22] H. Yang, X.P. Qin, J. Zhang, S.W. Wang, J. Ma, L.X. Wang, Q.T. Zhang, J. Alloys Compd. 509 (2011) 5274–5279. [23] J. Hosta^sa, L. Esposito, A. Piancastelli, J. Eur. Ceram. Soc. 32 (2012) 2949–2956. [24] J. Li, Q. Chen, G.Y. Feng, W.J. Wu, D.Q. Xiao, J.G. Zhu, Ceram. Inter. 38S (2012) S649–S652. [25] W. Zhang, T.C. Lu, N. Wei, Y.Z. Wang, B.M. Ma, F. Li, Z.W. Lu, J.Q. Qi, Opt. Mater. 34 (2012) 685–690. [26] W. Zhang, T.C. Lu, N. Wei, Y.Z. Wang, B.M. Ma, F. Li, Z.W. Lu, J.Q. Qi, J. Alloys Compd. 520 (2012) 36–41. [27] W. Zhang, T.C. Lu, N. Wei, B.Y. Ma, F. Li, Z.W. Lu, J.Q. Qi, Opt. Mater. 35 (2013) 2405–2410. [28] X.T. Chen, T.C. Lu, N. Wei, Z.W. Lu, W. Zhang, B.Y. Ma, Y.B. Guan, W. Liu, F.F. Jiang, J. Alloys Compd. 589 (2014) 448–454. [29] N. Wei, T.C. Lu, F. Li, W. Zhang, B.Y. Ma, J.Q. Qi, Appl. Phys. Lett. 101 (2012) 061902. [30] M.I. Mendelson, J. Am. Ceram. Soc. 52 (1969) 443–446. [31] Y.G. Wang, L.G. Zhang, Y. Fan, J.S. Luo, D.E. McCready, C.M. Wang, L.A. An, J. Am. Ceram. Soc. 88 (2005) 284–286. [32] S.H. Tong, T.C. Lu, W. Guo, Mater. Lett. 61 (2007) 4287–4289. [33] J.G. Li, T. Ikegami, J.H. Lee, T. Mori, Y. Yajima, J. Eur. Ceram. Soc. 20 (2000) 2395–2405. [34] K. Miyauchi, G. Toda, Opto-Ceram. Gihodo Syutsupan (1984) 49–63. [35] A. Ikesue, K. Yoshida, T. Yamamoto, I. Yamaga, J. Am. Ceram. Soc. 80 (6) (1996) 1517–1522. [36] A. Krell, J. Klimke, T. Hutzler, Opt. Mater. 31 (2009) 1144–1150. [37] X.W. Ba, J. Li, Y.P. Zeng, Y.B. Pan, B.X. Jiang, W.B. Liu, L. Wang, J. Liu, J.K. Guo, Ceram. Inter. 39 (4) (2013) 4639–4643.