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Intense visible upconversion emission in transparent (Ho3+, Er3+)-α-Sialon ceramics under 980 nm laser excitation
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Intense visible upconversion emission in transparent (Ho3+ , Er3+ )-␣-Sialon ceramics under 980 nm laser excitation Yuwaraj K. Kshetri a,∗ , Bhupendra Joshi b , Soo Wohn Lee b,∗ a b
Department of Advanced Materials Engineering, Sun Moon University, Chungnam 31460, Republic of Korea Department of Environment and Bio-Chemical Engineering, Sun Moon University, Chungnam 31460, Republic of Korea
a r t i c l e
i n f o
Article history: Received 11 May 2016 Received in revised form 27 June 2016 Accepted 2 July 2016 Available online xxx Keywords: Sialon Hot pressing Transparent ceramic Upconversion Phonon
a b s t r a c t A very intense infrared to visible frequency upconversion and near infrared frequency downconversion photoluminescence has been reported for the first time in Ho3+ and Er3+ co-doped transparent ␣-Sialon ceramics under 980 nm laser excitation. The ␣-Sialon ceramics were prepared by hot press sintering technique. Distinct upconversion bands at 553 nm, 656 nm, 671 nm and a downconversion band at 1530 nm were assigned to the transitions from 4 S3/2 to 4 I15/2 (Er3+ ), 5 F5 to 5 I8 (Ho3+ ), 4 F9/2 to 4 I15/2 (Er3+ ), and 4 I13/2 to 4 I15/2 (Er3+ ) respectively. The results show that there is very efficient energy transfer between Er3+ and Ho3+ ions. The quadratic dependence of upconversion intensity on the excitation power indicates that the upconversion process is governed by two photon absorption process. The highest phonon energy was found to be 855 cm−1 . Excellent up and down conversion emission properties together with moderately low phonon energy of the ␣-Sialon ceramics make the material potential candidate for multifunctional applications. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction Sialon ceramics are Si3 N4 based ceramics. Two major phases ␣ and -Sialon are isostructural with ␣ and -Si3 N4 and are represented by general formula Mx Si12 − m − n Alm + n On N16 − n (x = m/v, where x and v are the solubility and valency of the M ion, respectively) and Si6 − z Alz Oz N8 − z , respectively. In ␣-Sialon formation, the stabilizing cation M such as Li+ , Ca2+ , Mg2+ , and one of the most rare earth cations enters into two large interstices of the ␣-Si3 N4 (space group P31c) unit cell structure and m + n (Si N) bonds are replaced by m (Al N) and n (Al O) bonds in the crystal structure [1,2]. Metal cations used for the stabilization of the ␣-Si3 N4 are also responsible for the densification during the sintering, improvement of the microstructural characteristics and mechanical properties of these ceramics materials [3]. Research on ␣-Sialon ceramics has been concentrated mostly on stressed structural applications because of their outstanding thermo-mechanical properties such as high strength, superior wear resistance, excellent thermal shock resistance and low expansion coefficients [4–6]. But very little attention has been paid for functional (optical) applications. Karunaratne et al. [7] first
∗ Corresponding authors. E-mail addresses: [email protected] (Y.K. Kshetri), [email protected] (S.W. Lee).
reported a colored ␣- Sialon ceramics with relatively high optical transparency. Since then there have been several reports on translucent/transparent ␣-Sialon ceramics [5,8–11]. In these reports, ␣-Sialon ceramics were found to have better optical transparency because of their equiaxed grain morphology as compared to elongated grains of -Sialon. With improved transparency in both visible and infrared regions, the quest for photoluminescent ␣-Sialon ceramics began. Van Krevel et al. [12] first prepared and investigated luminescence properties of Tb3+ , Ce3+ and Eu2+ doped ␣-Sialon ceramics and reported low-energy 4f-5d transitions as compared to the luminescence of these ions doped in oxide host-lattices. Some luminescence properties on various rare earths doped ␣-Sialon ceramics have since been reported [13–17]. However, these works are mainly focused on downconversion luminescence (red shift) that results because of the change in oxidation state of the stabilizing cation under ultraviolet and visible excitations. There is almost no report on near-infrared to visible upconversion photoluminescence in ␣-Sialon ceramics so far. Upconversion materials have been proposed for several applications, such as lasers, solar cells, waveguides and display devices [18]. Most of the upconversion transparent materials are based on oxides, halides, and chalcogenides glasses. Heavy halides like chlorides, bromides, and iodides generally exhibit low phonon energies [19]. However, these are hygroscopic and suffer from poor thermal, chemical and mechanical stability and, therefore, are of limited use.
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Table 1 Sample codes and corresponding Ho and Er concentration. Sample Codes
Ho-concentration (mol%)
Er-concentration (mol%)
HE05 HE10 HE15 HE20 HE25 H30 E30a
0.5 1.0 1.5 2.0 2.5 3.0 0.0
2.5 2.0 1.5 1.0 0.5 0.0 3.0
a Photoluminescence of Er singly doped Sialon has been reported in our previous work (Ref. [31]). In the current study, the sample E30 has been prepared only for comparison of its absorption and photoluminescence spectra with other samples.
Oxides exhibit better chemical stability, but their phonon energies are relatively higher (more than 500 cm−1 ) [13,18,20]. For efficient upconversion process, the crystal structure and optical property of host material play important roles. Ideal host material should have low phonon energy, close lattice matches to dopant ions, adequate transparency within the wavelength range of interest, high optical damage threshold, and good chemical stability [21]. In this respect ␣-Sialon ceramic can be a potential host candidate for upconversion owing to its excellent thermo-mechanical properties. An additional advantage is that the ␣-Sialon family has the flexibility of varying chemical composition over a very wide range while maintaining the crystal structure, which favors the optimization and adjustment of photoluminescence properties [14]. While most lanthanide ions (Ln3+ ) can principally be expected to undergo upconversion process, the greatest upconversion efficiency under low excitation pump intensities at 980 nm is realized only for a few Ln3+ ions namely Erbium (Er3+ ), Holmium (Ho3+ ) and Thulium (Tm3+ ) [22,23]. There have been some reports of NIR to visible upconversion in Ho3+ and Er3+ doped glasses and ceramics [24–29]. In this paper, we report a very intense infrared to visible upconversion and downconversion photoluminescence emissions for the first time in Ho3+ and Er3+ co-doped transparent ␣-Sialon ceramics under 980 nm laser excitation. Phase and microstructure analysis has been done in relation to the XRD and TEM measurements. Spectroscopic properties in visible and infrared regions have been investigated and details of the energy transfer process between Er3+ and Ho3+ ions has been discussed. We also measured the phonon energy of ␣-Sialon ceramics from Raman spectroscopy. 2. Experimental Holmium (Ho) and Erbium (Er) singly doped and co-doped compositions were designed according to the ␣-Silaon formula Mx Si12− m+n Alm+n On N16−n ; x = m/3. The parameters m = n = 1.1 were chosen so as to get the final product within the ␣-Silaon phase area [1,8]. Sample codes and corresponding Ho and Er concentrations have been given in Table 1. The starting materials were high purity ␣-Si3 N4 powder (SN-E10, UBE Co., Japan), AlN (Grade F, Tokuyama Co., Japan), Al2 O3 (High purity chemicals Co., Ltd., Japan), Ho2 O3 (Sigma-Aldrich Co., China), and Er2 O3 (High purity chemicals Co LTD., Japan). These powders were mixed in appropriate proportion and ball milled with high purity Si3 N4 balls in absolute ethanol medium for 12 h. The ethanol was then completely removed by drying the milled slurry at 70 ◦ C for 4 h in a rotary evaporator. The dried mixture was then dry-milled with high purity Si3 N4 balls in air medium inside a sealed polyethylene jar for 12 h and finally sieved through 150 m mesh. The mixture of powders was packed into a boron nitride coated carbon mold and sintered in a hot press furnace (MVHP, Monocerapia Co. Ltd., Korea) at 1800◦ C and 25 MPa uniaxial pressures in nitrogen environment for 2 h. The sintered samples in the form of circular plates each of diameter
Fig. 1. XRD patterns of different ␣-Sialon samples. The pattern at bottom axis is the standard pattern of JCPDS 33-261 of ␣-Sialon. Table 2 Unit cell parameters of Ho singly doped and Ho, Er co-doped ␣-Sialon samples. Samples
˙ a (A)
˙ c (A)
c/a
Volume (A˙ 3 )
HE05 HE10 HE15 HE20 HE25 H30
7.8037 7.8075 7.8121 7.8038 7.8073 7.8108
5.6630 5.6823 5.6751 5.6629 5.6649 5.6840
0.726 0.728 0.726 0.726 0.726 0.728
298.66 299.97 299.95 299.67 299.04 300.31
50 mm were ground, cut and polished for various mechanical and spectroscopic measurements. Density was measured by Archimedes immersion method in distilled water. Hardness and fracture toughness were measured using a Vickers diamond indenter with 98 N load. Phase of the specimens was analyzed by X-ray diffractometer (XRD; Rigaku D/Max 2200HR diffractometer, Japan) using Cu K␣ radiation over a 2 range of 10◦ to 90◦ at a scanning rate of 1◦ per minute. Microstructure and elemental analysis were carried out by transmission electron microscope (TEM; JEM-4010, JEOL, Japan) operated at 400 kV. Fracture surface morphology was observed by scanning electron microscope (SEM; SNE-3000, SEC, Korea). Optical absorption and transmission spectra of 0.50 mm thick samples with polished surfaces were measured by UV/Vis/NIR Jasco 570 Spectrometer as well as FTIR Vertex 70, Bruker spectrometer. Photoluminescence (PL) spectra under 980 nm continuous wave laser excitation were recorded by Shamrock Spectrograph (Shamrock 303i, Andor). Raman spectra were measured in backscattering configuration using Raman spectrometer (LabRam HR, Horiba) coupled with 532 nm green laser. All the measurements were carried out at room temperature. 3. Results and discussion 3.1. Phase, microstructure analysis and mechanical properties The XRD patterns of Ho singly doped and Ho, Er co-doped samples are shown in Fig. 1. All the sintered samples contain near ␣-Sialon phase with a trace amount of -Sialon phase. We observed that all the diffraction peaks were in good agreement with those of the standard JCPDS file 33-261 of ␣-Sialon as shown in Fig. 1. The lattice parameters presented in Table 2 were calculated from corre-
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Fig. 2. SEM micrograph of fractured surfaces of different ␣-Sialon samples; (a) HE05, (b) HE10, (c) HE15 (d) HE20 (e) HE25 and (f) H30.
Fig. 3. (a) TEM micrograph of H30; Ho singly doped ␣-Sialon. (b) Selected area diffraction (SAD) pattern of region DP in (a). (c) A triple junction. EDS spectra 1, 2 and 3 correspond to the points 1, 2 and 3 respectively in (a).
sponding diffraction patterns of the samples. Silicon standard peaks were used to calculate the lattice parameters. The lattice parameters and unit cell volume of all the samples are within the expected range of ␣-Sialon unit cell [30]. Although there is a slight variation in the values of a and c of the individual sample, the ratio c/a is fairly constant (0.726–0.728) for all the samples. SEM images of the frac-
ture surface in Fig. 2 show that all the samples are composed of isometric, equiaxed polyhedral grains which are the most common ␣-Sialon grain morphology [1]. Such grain morphology gives better optical transparency [7,8]. The fracture modes are mainly intercrystalline. In order to observe the effect of co-doping, we carried out an in-depth TEM analysis of Ho singly doped and Ho, Er co-doped
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Fig. 4. (a) TEM micrograph of HE10; Ho and Er co-doped ␣-Sialon. (b) Selected area diffraction (SAD) pattern of region DP in (a). (c) A triple junction. EDS spectra 1, 2, 3, 4 and 5 correspond to the points 1, 2, 3, 4 and 5 respectively in (a).
samples. Fig. 3 shows the TEM micrographs along with EDS spectra at selected points of sample H30. Typical ␣-Sialon grain morphology of H30 can be seen in Fig. 3(a). In addition, the EDS spectra 1 and 2 of the points inside the grains show the presence of Ho3+ ions in the grains indicating that Ho3+ ions have stabilized the ␣-Sialon. This is also confirmed by the SAD pattern in Fig. 3(b). The crystal planes (100) and (102) in the SAD patterns also belong to the ␣Sialon phase of the XRD pattern of H30 as shown in Fig. 1. Similar results can be seen in Fig. 4. for the Ho, Er co-doped sample HE10. The crystal planes (100) and (101) in the SAD patterns in Fig. 4(b) are also consistent with the XRD pattern of HE10 indicating the ␣Sialon phase. EDS spectra 1, 2 and 3 of the points within the grains show the presence of both Ho3+ and Er3+ ions inside the grains. Conformation of the ␣-Sialon phase via XRD and SAD pattern analysis in conjunction with the result of EDS analysis of the presence of both Ho3+ and Er3+ ions inside the ␣-Sialon grains indicate that both Ho3+ and Er3+ ions have been successfully doped into the interstitial sites of the ␣-Si3 N4 matrix to form duel ions stabilized (Ho3+ , Er3+ )-␣-Sialon. The co-existence of Ho3+ and Er3+ ions into the lattice is very important for the effective energy transfer between the ions. Some Ho3+ and Er3+ ions were observed also at triple junctions (point 4) and grain boundaries (point 5) as indicated by the corresponding EDS spectra. In our previous works [31,32], all the Er3+ ions were found at grain boundaries and triple junctions while only smaller co-doped cations like Mg2+ or Yb3+ were found inside the grains.
The high-resolution transmission electron microscopy (HRTEM) image of H30 is shown in Fig. 5. A thin intergranular grain boundary phase of thickness 2.15 nm can be seen in Fig. 5(c). The calculated interplanar spacing of 0.25 nm in Fig. 5(d) corresponds to (210) plane which is the strongest ␣-Sialon peak in the XRD spectrum. The d-spacing of (101) plane in SAD pattern of Fig. 5(e) was found to be 0.44 nm identical to that in Fig. 5(b). Similarly, HRTEM image of HE10 is shown in Fig. 6. The intergranular grain boundary phase is reduced approximately to 1.50 nm in HE10, as shown in Fig. 6(c), as compared to that of the Ho singly doped sample, H30. Hence, because of co-doping of Ho3+ and Er3+ ions the amount of intergranular phase is reduced by the transient liquid phase being absorbed into the matrix of ␣-Sialon phase during sintering [5]. The interplanar spacing of 0.28 nm and 0.68 nm in Fig. 6(b) and (d) correspond respectively to the planes (002) and (100) in the XRD spectrum. We observed a very good agreement between all the calculated dspacing of the indicated lattice planes in HRTEM images of both H30 and HE10 samples with the d values of the corresponding planes of XRD spectra of JCPDS 33-261 of ␣-Sialon. Hardness, fracture toughness and density of Ho singly doped and Ho, Er co-doped samples are tabulated in Table 3. The inherent outstanding mechanical properties of ␣-Sialon ceramics are consistently retained by all the sintered samples with hardness above 20 GPa and fracture toughness above 5 MPa m1/2 . The ␣-Sialon ceramics exhibit high hardness due to their small grain size and high atomic density in unit cell structure while thin glassy phase
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Fig. 5. (a) HRTEM micrograph of H30; (b), (c) and (d) are the magnified images of the selected areas in (a); (e) is the selected area diffraction (SAD) pattern of the area (b).
Fig. 6. (a) HRTEM micrograph of HE10; (b), (c) and (d) are the magnified images of the selected areas in (a); (e) is the selected area diffraction (SAD) pattern of the area (d).
Table 3 Mechanical properties and densities of different ␣-Sialon samples.
3.2. Transmission and absorption spectra
Samples
Hardness (GPa)
Fracture toughness (MPa m1/2 )
Density (g/cm3 )
HE05 HE10 HE15 HE20 HE25 H30
20.24 20.54 20.45 20.35 20.19 20.33
5.18 5.59 5.44 5.61 5.52 5.59
3.36 3.38 3.41 3.42 3.42 3.43
that exists as intergranular phase provides high fracture toughness. The observed hardness and fracture toughness of the sintered ␣Sialon ceramics are better than that for yttrium aluminum garnet (YAG), Y2 O3 , and other optically active polycrystalline transparent ceramics available commercially [31]. Measured low densities of the samples additionally make ␣-Sialon ceramics a light weight material.
Transmission spectra of Ho singly doped and Ho, Er co-doped samples each of thickness 0.50 mm are shown in Fig. 7. The transmission measurements were performed in the wavelength range from 200 nm to 5500 nm in order to cover visible and IR regions. The transparency first increases with the increasing wavelength, reaches as high as 80% in the mid-IR region and then decreases. Zhong et al. [9] performed the first principle calculation to understand the underlying transparency mechanism in ␣-Sialon. According to them, the doped rare-earth atom interacts strongly with the doping states of ␣-Sialon, resulting in the increase in the optical gap, which suppresses the photoabsorption in the wavelength region between 1000–4500 nm and leads to the better optical transparency. The optical image of the samples in the inset shows that the sintered samples have very good transparency in visible region as well. However, Ho singly doped sample i.e., H30 has relatively poor transmittance as compared to the Ho, Er co-
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Fig. 7. Transmission spectra of all the ␣-Sialon samples of thickness 0.50 mm. Inset is the optical image of the samples.
Fig. 9. Photoluminescence spectra of different ␣-Sialon samples under 980 nm excitation. Inset (a) is the magnified view of the upconversion emission band at 766 nm, inset (b) is the down conversion spectra. The left inset is the photograph of green and red luminescence from E30 and HE05 respectively as observed by the naked eye. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
Fig. 8. Absorption spectra of Ho singly doped, Er singly doped and Ho, Er co-doped ␣-Sialon samples.
doped samples. It can be inferred that the co-doping of Ho and Er has significantly improved the optical transparency. Improvement in transparency can be attributed to the following reasons. Firstly, co-doping of Ho and Er has yielded better grain morphology as compared to the Ho single doping which is evident from Fig. 2. Secondly, the intergranular grain boundary phase, which is mostly responsible for the loss due to scattering in polycrystalline ceramics [8,10], has been reduced in the co-doped sample as is evident from Figs. 5 and 6. The reduction in the grain boundary phase obviously offers a more smooth path for light propagation across the grain boundaries. Improved optical transparency could also be due to the reduced residual pores, which act as light scattering centers, in the sintered body of the co-doped samples as compared to that in Ho3+ singly doped sample. Absorption spectra of Ho singly doped, Er singly doped and Ho, Er co-doped samples in the wavelength range of 200 nm to 2500 nm are shown in Fig. 8. All the characteristics absorption transitions of Ho3+ and Er3+ ions from their corresponding ground states 5 I8 and 4 I15/2 have been thoroughly labeled. In the absorption spectrum of H30, there is no absorption band around 980 nm. Hence,
Fig. 10. Er concentration dependence of intensities at 553, 656 and 671 nm emissions in ␣-Sialon samples.
conventional 980 nm pumping cannot be used for the Ho singly doped ␣- Sialon. However, the presence of absorption band 4 I11/2 of Er3+ at 980 nm in the HE10 indicates that the co-doped samples can be excited by 980 nm laser. Moreover, it can be seen that a pair of absorption bands, in the visible region, of Er3+ : 4 I15/2 → 4 S3/2 and 4 I15/2 → 4 F9/2 overlap that of Ho3+ :5 I8 → 5 S2 (5 F4 ) and 5 I8 → 5 F5 transitions, respectively. Hence, a sensitized upconversion luminescence [33,34] can be expected in Ho, Er co-doped ␣-Sialon. 3.3. Photoluminescence properties The room temperature photoluminescence spectra of all the samples under 980 nm excitation have been shown in Fig. 9. No upconversion emission can be seen for Ho singly doped sample, H30, because Ho3+ cannot be excited at 980 nm. For Er singly
Please cite this article in press as: Y.K. Kshetri, et al., Intense visible upconversion emission in transparent (Ho3+ , Er3+ )-␣-Sialon ceramics under 980 nm laser excitation, J Eur Ceram Soc (2016), http://dx.doi.org/10.1016/j.jeurceramsoc.2016.07.005
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Fig. 11. Power dependence and the photonic process of upconversion in Ho and Er co-doped ␣-Sialon samples.
doped sample, E30, there is a strong green emission and week red and near infrared emissions. A frequency downconversion band is also observed around 1530 nm as shown in the inset (b) of Fig. 9. The emission spectra drastically change in Ho and Er codoped samples. A very intense red emission can be seen around 671 nm along with a new weak emission at 766 nm in all the co-doped samples. However, the intensities of green and NIR emissions decrease in the co-doped samples as compared to that in E30. The emissions at 535 nm, 553 nm, 671 nm, 766 nm, 812 nm and 1530 nm are assigned to Er3+ :2 H11/2 → 4 I15/2 , Er3+ :4 S3/2 → 4 I15/2 , Er3+ : 4 F9/2 → 4 I15/2 , Ho3+ : 5 S2 (5 F4 ) → 5 I7 , Er3+ :4 I9/2 → 4 I15/2 , and Er3+ :4 I13/2 → 4 I15/2 respectively. It is found that when Er concentration is increased from zero to 2.5 mol%, a new peak gradually appears at 656 nm whose intensity also increases while the intensity at 671 nm emission first increases up to 1.5 mol% and then decreases. This behavior has been depicted in Fig. 10. The bottom axis of the figure labels the Er concentration and the top axis represents the corresponding sample. We believe that at higher Er concentration energy transfer from Er3+ to Ho3+ becomes more efficient thereby increasing the intensity at 656 nm whereas the emission at 671 nm from Er3+ : 4 F9/2 → 4 I15/2 transition itself is reduced. Hence, the red emission at 656 nm can be assigned to Ho3+ : 5 F → 5 I transition. The emissions at 656 nm and 766 nm are the 5 8 strong evidence of the energy transfer from Er3+ to Ho3+ . It should be emphasized here that the green emission in E30 and red emissions in all co-doped samples are intense enough to be observed even by the naked eye at excitation power as low as 75 mW (see Fig. 9). The first reason for the observed efficient upconversion is the co-existence of Er3+ and Ho3+ ions in the lattice of ␣-Sialon. As revealed earlier in EDS analysis, Er3+ and Ho3+ ions have been successfully doped into the two large interstices of the ␣-Sialon. This makes the ions sufficiently close for effective energy transfer from Er3+ to Ho3+ . The energy transfer probability is determined by the overlapping of the wave functions of the sensitizer and activator and this probability is high only when activator and sensitizer occupy adjacent lattice sites [33]. The second reason for the efficient
Fig. 12. Simplified energy level diagram of Er3+ and Ho3+ and involved energy transfer mechanisms.
upconversion can be ascribed directly to the different crystal-fields around the trivalent lanthanide ions in matrices of ␣-Sialon. The unit cell of ␣-Sialon has hexagonal geometry and can, therefore, be considered a low symmetry host. It is known that the low symmetry hosts typically exert a crystal field containing more uneven components around the dopant ions compared to high symmetry counterparts [20]. The uneven components in ␣-Sialon matrix therefore enhance the electronic coupling between 4f energy levels and higher electronic configuration and subsequently increase f-f transition probabilities of the dopant ions giving enhanced upconversion emission. Moreover, decrease in cationic size or unit cell volume of the host can cause an increase in the crystal field strength around the dopant ions and lead to enhanced upconversion effi-
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Fig. 13. CIE chromaticity diagram of Er singly doped and Ho, Er co-doped ␣-Sialon ceramics. The top inset shows the color coordinates and the color purity of the samples.
ciency. This fact is consistent with our observation that HE05 has the strongest upconversion emission which has the least unit cell volume. The pump power dependence of luminescence intensities has been investigated to understand the upconversion mechanism. A typical behavior of HE05 for the incident power from 75 mW to 300 mW has been shown in Fig. 11(a). It has been shown that for the case of a small upconversion rates, the luminescence intensity I is proportional to the nth power of the absorbed pump power P. Mathematically [35], I ∝ P n ⇒ Log(I) ∝ nLog(P)
(1)
where n, called the order of the upconversion process, is the number of pump photons absorbed per upconverted photon and is given by the slope of the luminescence intensity versus pump power in double-logarithmic representation. The results presented in Fig. 11(b)–(f) within the above mentioned power range for all the samples show that the value of n is nearly 2. Hence, the luminescence intensity shows the quadratic dependence on the pump power indicating a two photon absorption process predominantly populates green and red emitting states. Some deviation of n from typical values can be attributed to the competition between the linear decay and the upconversion processes for the depletion of the intermediate excited states [35].
Based on the simplified energy level diagram, Fig. 12, a mechanism for the observed photoluminescence phenomena has been presented. Excited state absorption (ESA) and energy transfer (ET) upconversion are the dominant upconversion mechanisms in Er3+ , Ho3+ doped systems [28,36,37]. We first discuss the ESA mechanism of Er3+ excitation. Upon 980 nm excitation, Er3+ ions are first excited from 4 I15/2 to 4 I11/2 by the ground state absorption (GSA). On one hand, further excitation of Er3+ : 4 I11/2 takes place by excited state absorption ESA1 to make the Er3+ : 4 F7/2 level populated. The populated Er3+ : 4 F7/2 level then relaxes rapidly and non-radiatively to the next lower levels 2 H11/2 and 4 S3/2 . The final radiative transitions 2 H11/2 → 4 I15/2 and 4 S3/2 → 4 I15/2 produce green emissions at 535 nm and 553 nm respectively. On the other hand, some of Er3+ ions in 4 I11/2 nonradiatively decay to 4 I13/2 . And a subsequent radiative transition to 4 I15/2 yields 1530 nm downconversion emission. Moreover, some of the Er3+ ions in 4 I13/2 can also be excited to 4 F9/2 via ESA2 to produce either red emission at 671 nm by 4 F9/2 → 4 I15/2 transition or NIR emission at 812 nm by 4 I9/2 → 4 I15/2 after a nonradiative relaxation of the 4 F9/2 level. We now consider the energy transfer (ET) from Er3+ to Ho3+ . Here Er3+ is a sensitizer and Ho3+ is an activator. The excited Er3+ ions in 4 I11/2 state may transfer their energy to Ho3+ ions. The 5 I6 level of Ho3+ lies about 1290 cm−1 below the 4 I11/2 level of Er3+ . This energy can easily be taken up [28] by the ␣-Sialon lattice since its vibrational mode exists at 760 cm−1 . Therefore Ho3+ ions
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in the ground state 5 I8 are excited to the 5 I6 state by the neighboring Er3+ ions via ET1: 4 I11/2 Er3+ + 5 I8 Ho3+ → 4 I15/2 Er3+ + 5 I6 Ho3+ . The nonradiative relaxations of 5 I6 → 5 I7 populate 5 I7 level of the Ho3+ ions. The second step involves the excitation of Ho3+ either from 5 I6 to 5 S2 , 5 F4 state via ET2: 4 I11/2 Er3+ + 5 I6 Ho3+ → 4 I15/2 Er3+ + 5 S2 , 5 F4 Ho3+ , or from 5 I7 to 5 F5 state via a phonon-assisted ET3: 4 I11/2 Er3+ + 5 I7 Ho3+ → 4 I15/2 Er3+ + 5 F5 Ho3+ . Moreover, 5 F5 can also be populated by the nonradiative decay of 5 S2 , 5 F4 states of Ho3+ . The radiative decay of Ho3+ : 5 F5 → 5 I8 gives red emissions at 656 nm while the decay of Ho3+ : 5 S2 , 5 F4 → 5 I7 gives NIR emission at 766 nm. Because of the multiphonon relaxation, Ho3+ ions in 5 I6 are more likely to relax to lower level 5 I7 and those at 5 S2 level are more likely to relax to lower 5 F5 level nonradiatively [38]. These two processes decrease the population of 5 S2 level and increase that of 5 F5 . Therefore, the red emission is dominant and NIR emission is weak in Ho3+ . At higher Er3+ concentration, yet another energy transfer process ET4 becomes significant [24]: 4 F9/2 Er3+ + 5 I7 Ho3+ → 4 I13/2 Er3+ + 5 F5 Ho3+ . As a consequence, the red emission intensity at 656 nm due to Ho3+ : 5 F5 → 5 I8 increases while that at 671 nm due to Er3+ :4 F9/2 → 4 I15/2 decreases which were observed in the PL emission spectra of Fig. 9. Fig. 13 is the Commission Internationale de L’Eclairage (CIE) chromaticity diagram of all the samples under 980 nm excitation. The resultant CIE coordinates of the samples listed in the inset were calculated from the corresponding photoluminescence emission spectra of Fig. 9. The color purity was calculated using the method discussed in our previous work [31]. It can be seen that Er single doped sample falls in the green region while all the Ho and Er co-doped samples fall in the red region. The high color purity of the samples broadens the fields of application of this material as upconversion phosphors.
3.4. Raman spectra Raman spectra of the sintered samples are shown in Fig. 14. The sample H30 shows a broad Raman band centered around 650 cm−1 . The inset shows the Gaussian resolved components of the sample H30 with the highest vibrational band at 800 cm−1 . The fundamental structural units in Si3 N4 crystal and Sialons are the Si–N4 tetrahedra, linked in such a way that the N– Si3 triangular unit is approximately planar. Therefore, the vibrational spectra of Sialons usually display a broad band attesting to the significantly distorted structure of these Si– N4 tetrahedra and N–Si3 triangular units [39]. Spectra of Ho, Er co-doped samples are significantly different from that of Ho single doped sample. It is evident that the incorporation of Er3+ ions splits the peaks into two distinct low and high frequency bands below and above 650 cm−1 . First principle calculations [40–42] of vibrational density of states of ␣-Si3 N4 have also shown that two different modes of vibration of N atoms give rise to the well separated bands below and above 600 cm−1 . Si N bond is much stiffer against stretching than bending [42]. Low and high frequency phonon bands below and above 650 cm−1 in our case can, therefore, be ascribed respectively to bond bending (rocking motion) and bond stretching vibrations of N atoms in N– Si3 triangular units. The very low frequency band at 135 cm−1 is attributed to the vibrations of the acoustic and optic branches distorting the whole Si N bonding network. The observed vibrational frequencies of 468, 515, 545, 760 and 855 cm−1 are in good agreement with other reports [43,44] of experimental and calculational results for ␣-Si3 N4 . The highest phonon energy of 855 cm−1 for Ho, Er co-doped ␣Sialon is still less than that of the well known upconversion and laser host materials such as phosphate glass (1200 cm−1 ), Silica glass (1100 cm−1 ), LaPO4 (1050 cm−1 ), YAG (860 cm−1 ), and fluorophosphate glass (1060 cm−1 ) [18,26]. The moderately low value
9
Fig. 14. Raman Spectra of different ␣-Sialon samples. The inset shows the Gaussian resolved components of Raman spectrum of H30 sample.
of phonon energy has the effect of reducing the non-radiative multiphonon relaxation in the ␣-Sialon host and enhancing the novel energy transfer process between Er3+ and Ho3+ ions for the efficient photoluminescence. 4. Conclusions We have demonstrated a very intense infrared to visible frequency upconversion and near infrared frequency downconversion photoluminescence for the first time in Ho3+ and Er3+ co-doped transparent ␣-Sialon ceramics, sintered by hot press technique, under 980 nm laser excitation. The intense upconversion emission was due to the efficient energy transfer process between Er3+ and Ho3+ ions as a result of co-existence of the ions in the interstices of the ␣-Sialon lattice. Formation of the ␣-Sialon phase and existence of both Er3+ and Ho3+ ions in the lattice were confirmed by XRD, TEM, and EDS analysis. Transparency of the sintered samples was improved by the co-doping of Er3+ and Ho3+ ions due to the reduced intergranular boundary phase in the co-doped samples. Raman spectroscopic study revealed that the ␣-Sialon ceramics have moderately low phonon energy as compared to some well known phosphate and oxide glasses and ceramics. The inherent outstanding thermo-chemical and mechanical stability along with the up and downconversion emission properties and moderately low phonon energy of (Ho3+ , Er3+ )-␣-Sialon ceramics make the material promising candidate for multifunctional applications. Acknowledgment This research was supported by Global Research Laboratory (GRL) program of the National Research Foundation of Korea (NRF) funded by the Ministry of Education, Science and Technology (MEST), Republic of Korea grant number (2010-00339). References [1] G.Z. Cao, R. Metselaar, ␣’-Sialon ceramics: a review, Chem. Mater. 3 (1991) 242–252. [2] S. Hampshire, H.K. Park, D.P. Thompson, K.H. Jack, ␣’-Sialon ceramics, Nature 274 (1978) 880–882. [3] C. Santos, K. Strecker, P.A. Suzuki, S. Kycia, O.M.M. Silva, C.R.M. Silva, Stabilization of ␣-SiAlONs using a rare-earth mixed oxide (RE2 O3 ) as sintering additive, Mater. Res. Bull. 40 (2005) 1094–1103.
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Intense visible upconversion emission in transparent (Ho3+ , Er3+ )-␣-Sialon ceramics under 980 nm laser excitation Yuwaraj K. Kshetri a,∗ , Bhupendra Joshi b , Soo Wohn Lee b,∗ a b
Department of Advanced Materials Engineering, Sun Moon University, Chungnam 31460, Republic of Korea Department of Environment and Bio-Chemical Engineering, Sun Moon University, Chungnam 31460, Republic of Korea
a r t i c l e
i n f o
Article history: Received 11 May 2016 Received in revised form 27 June 2016 Accepted 2 July 2016 Available online xxx Keywords: Sialon Hot pressing Transparent ceramic Upconversion Phonon
a b s t r a c t A very intense infrared to visible frequency upconversion and near infrared frequency downconversion photoluminescence has been reported for the first time in Ho3+ and Er3+ co-doped transparent ␣-Sialon ceramics under 980 nm laser excitation. The ␣-Sialon ceramics were prepared by hot press sintering technique. Distinct upconversion bands at 553 nm, 656 nm, 671 nm and a downconversion band at 1530 nm were assigned to the transitions from 4 S3/2 to 4 I15/2 (Er3+ ), 5 F5 to 5 I8 (Ho3+ ), 4 F9/2 to 4 I15/2 (Er3+ ), and 4 I13/2 to 4 I15/2 (Er3+ ) respectively. The results show that there is very efficient energy transfer between Er3+ and Ho3+ ions. The quadratic dependence of upconversion intensity on the excitation power indicates that the upconversion process is governed by two photon absorption process. The highest phonon energy was found to be 855 cm−1 . Excellent up and down conversion emission properties together with moderately low phonon energy of the ␣-Sialon ceramics make the material potential candidate for multifunctional applications. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction Sialon ceramics are Si3 N4 based ceramics. Two major phases ␣ and -Sialon are isostructural with ␣ and -Si3 N4 and are represented by general formula Mx Si12 − m − n Alm + n On N16 − n (x = m/v, where x and v are the solubility and valency of the M ion, respectively) and Si6 − z Alz Oz N8 − z , respectively. In ␣-Sialon formation, the stabilizing cation M such as Li+ , Ca2+ , Mg2+ , and one of the most rare earth cations enters into two large interstices of the ␣-Si3 N4 (space group P31c) unit cell structure and m + n (Si N) bonds are replaced by m (Al N) and n (Al O) bonds in the crystal structure [1,2]. Metal cations used for the stabilization of the ␣-Si3 N4 are also responsible for the densification during the sintering, improvement of the microstructural characteristics and mechanical properties of these ceramics materials [3]. Research on ␣-Sialon ceramics has been concentrated mostly on stressed structural applications because of their outstanding thermo-mechanical properties such as high strength, superior wear resistance, excellent thermal shock resistance and low expansion coefficients [4–6]. But very little attention has been paid for functional (optical) applications. Karunaratne et al. [7] first
∗ Corresponding authors. E-mail addresses: [email protected] (Y.K. Kshetri), [email protected] (S.W. Lee).
reported a colored ␣- Sialon ceramics with relatively high optical transparency. Since then there have been several reports on translucent/transparent ␣-Sialon ceramics [5,8–11]. In these reports, ␣-Sialon ceramics were found to have better optical transparency because of their equiaxed grain morphology as compared to elongated grains of -Sialon. With improved transparency in both visible and infrared regions, the quest for photoluminescent ␣-Sialon ceramics began. Van Krevel et al. [12] first prepared and investigated luminescence properties of Tb3+ , Ce3+ and Eu2+ doped ␣-Sialon ceramics and reported low-energy 4f-5d transitions as compared to the luminescence of these ions doped in oxide host-lattices. Some luminescence properties on various rare earths doped ␣-Sialon ceramics have since been reported [13–17]. However, these works are mainly focused on downconversion luminescence (red shift) that results because of the change in oxidation state of the stabilizing cation under ultraviolet and visible excitations. There is almost no report on near-infrared to visible upconversion photoluminescence in ␣-Sialon ceramics so far. Upconversion materials have been proposed for several applications, such as lasers, solar cells, waveguides and display devices [18]. Most of the upconversion transparent materials are based on oxides, halides, and chalcogenides glasses. Heavy halides like chlorides, bromides, and iodides generally exhibit low phonon energies [19]. However, these are hygroscopic and suffer from poor thermal, chemical and mechanical stability and, therefore, are of limited use.
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Table 1 Sample codes and corresponding Ho and Er concentration. Sample Codes
Ho-concentration (mol%)
Er-concentration (mol%)
HE05 HE10 HE15 HE20 HE25 H30 E30a
0.5 1.0 1.5 2.0 2.5 3.0 0.0
2.5 2.0 1.5 1.0 0.5 0.0 3.0
a Photoluminescence of Er singly doped Sialon has been reported in our previous work (Ref. [31]). In the current study, the sample E30 has been prepared only for comparison of its absorption and photoluminescence spectra with other samples.
Oxides exhibit better chemical stability, but their phonon energies are relatively higher (more than 500 cm−1 ) [13,18,20]. For efficient upconversion process, the crystal structure and optical property of host material play important roles. Ideal host material should have low phonon energy, close lattice matches to dopant ions, adequate transparency within the wavelength range of interest, high optical damage threshold, and good chemical stability [21]. In this respect ␣-Sialon ceramic can be a potential host candidate for upconversion owing to its excellent thermo-mechanical properties. An additional advantage is that the ␣-Sialon family has the flexibility of varying chemical composition over a very wide range while maintaining the crystal structure, which favors the optimization and adjustment of photoluminescence properties [14]. While most lanthanide ions (Ln3+ ) can principally be expected to undergo upconversion process, the greatest upconversion efficiency under low excitation pump intensities at 980 nm is realized only for a few Ln3+ ions namely Erbium (Er3+ ), Holmium (Ho3+ ) and Thulium (Tm3+ ) [22,23]. There have been some reports of NIR to visible upconversion in Ho3+ and Er3+ doped glasses and ceramics [24–29]. In this paper, we report a very intense infrared to visible upconversion and downconversion photoluminescence emissions for the first time in Ho3+ and Er3+ co-doped transparent ␣-Sialon ceramics under 980 nm laser excitation. Phase and microstructure analysis has been done in relation to the XRD and TEM measurements. Spectroscopic properties in visible and infrared regions have been investigated and details of the energy transfer process between Er3+ and Ho3+ ions has been discussed. We also measured the phonon energy of ␣-Sialon ceramics from Raman spectroscopy. 2. Experimental Holmium (Ho) and Erbium (Er) singly doped and co-doped compositions were designed according to the ␣-Silaon formula Mx Si12− m+n Alm+n On N16−n ; x = m/3. The parameters m = n = 1.1 were chosen so as to get the final product within the ␣-Silaon phase area [1,8]. Sample codes and corresponding Ho and Er concentrations have been given in Table 1. The starting materials were high purity ␣-Si3 N4 powder (SN-E10, UBE Co., Japan), AlN (Grade F, Tokuyama Co., Japan), Al2 O3 (High purity chemicals Co., Ltd., Japan), Ho2 O3 (Sigma-Aldrich Co., China), and Er2 O3 (High purity chemicals Co LTD., Japan). These powders were mixed in appropriate proportion and ball milled with high purity Si3 N4 balls in absolute ethanol medium for 12 h. The ethanol was then completely removed by drying the milled slurry at 70 ◦ C for 4 h in a rotary evaporator. The dried mixture was then dry-milled with high purity Si3 N4 balls in air medium inside a sealed polyethylene jar for 12 h and finally sieved through 150 m mesh. The mixture of powders was packed into a boron nitride coated carbon mold and sintered in a hot press furnace (MVHP, Monocerapia Co. Ltd., Korea) at 1800◦ C and 25 MPa uniaxial pressures in nitrogen environment for 2 h. The sintered samples in the form of circular plates each of diameter
Fig. 1. XRD patterns of different ␣-Sialon samples. The pattern at bottom axis is the standard pattern of JCPDS 33-261 of ␣-Sialon. Table 2 Unit cell parameters of Ho singly doped and Ho, Er co-doped ␣-Sialon samples. Samples
˙ a (A)
˙ c (A)
c/a
Volume (A˙ 3 )
HE05 HE10 HE15 HE20 HE25 H30
7.8037 7.8075 7.8121 7.8038 7.8073 7.8108
5.6630 5.6823 5.6751 5.6629 5.6649 5.6840
0.726 0.728 0.726 0.726 0.726 0.728
298.66 299.97 299.95 299.67 299.04 300.31
50 mm were ground, cut and polished for various mechanical and spectroscopic measurements. Density was measured by Archimedes immersion method in distilled water. Hardness and fracture toughness were measured using a Vickers diamond indenter with 98 N load. Phase of the specimens was analyzed by X-ray diffractometer (XRD; Rigaku D/Max 2200HR diffractometer, Japan) using Cu K␣ radiation over a 2 range of 10◦ to 90◦ at a scanning rate of 1◦ per minute. Microstructure and elemental analysis were carried out by transmission electron microscope (TEM; JEM-4010, JEOL, Japan) operated at 400 kV. Fracture surface morphology was observed by scanning electron microscope (SEM; SNE-3000, SEC, Korea). Optical absorption and transmission spectra of 0.50 mm thick samples with polished surfaces were measured by UV/Vis/NIR Jasco 570 Spectrometer as well as FTIR Vertex 70, Bruker spectrometer. Photoluminescence (PL) spectra under 980 nm continuous wave laser excitation were recorded by Shamrock Spectrograph (Shamrock 303i, Andor). Raman spectra were measured in backscattering configuration using Raman spectrometer (LabRam HR, Horiba) coupled with 532 nm green laser. All the measurements were carried out at room temperature. 3. Results and discussion 3.1. Phase, microstructure analysis and mechanical properties The XRD patterns of Ho singly doped and Ho, Er co-doped samples are shown in Fig. 1. All the sintered samples contain near ␣-Sialon phase with a trace amount of -Sialon phase. We observed that all the diffraction peaks were in good agreement with those of the standard JCPDS file 33-261 of ␣-Sialon as shown in Fig. 1. The lattice parameters presented in Table 2 were calculated from corre-
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Fig. 2. SEM micrograph of fractured surfaces of different ␣-Sialon samples; (a) HE05, (b) HE10, (c) HE15 (d) HE20 (e) HE25 and (f) H30.
Fig. 3. (a) TEM micrograph of H30; Ho singly doped ␣-Sialon. (b) Selected area diffraction (SAD) pattern of region DP in (a). (c) A triple junction. EDS spectra 1, 2 and 3 correspond to the points 1, 2 and 3 respectively in (a).
sponding diffraction patterns of the samples. Silicon standard peaks were used to calculate the lattice parameters. The lattice parameters and unit cell volume of all the samples are within the expected range of ␣-Sialon unit cell [30]. Although there is a slight variation in the values of a and c of the individual sample, the ratio c/a is fairly constant (0.726–0.728) for all the samples. SEM images of the frac-
ture surface in Fig. 2 show that all the samples are composed of isometric, equiaxed polyhedral grains which are the most common ␣-Sialon grain morphology [1]. Such grain morphology gives better optical transparency [7,8]. The fracture modes are mainly intercrystalline. In order to observe the effect of co-doping, we carried out an in-depth TEM analysis of Ho singly doped and Ho, Er co-doped
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Fig. 4. (a) TEM micrograph of HE10; Ho and Er co-doped ␣-Sialon. (b) Selected area diffraction (SAD) pattern of region DP in (a). (c) A triple junction. EDS spectra 1, 2, 3, 4 and 5 correspond to the points 1, 2, 3, 4 and 5 respectively in (a).
samples. Fig. 3 shows the TEM micrographs along with EDS spectra at selected points of sample H30. Typical ␣-Sialon grain morphology of H30 can be seen in Fig. 3(a). In addition, the EDS spectra 1 and 2 of the points inside the grains show the presence of Ho3+ ions in the grains indicating that Ho3+ ions have stabilized the ␣-Sialon. This is also confirmed by the SAD pattern in Fig. 3(b). The crystal planes (100) and (102) in the SAD patterns also belong to the ␣Sialon phase of the XRD pattern of H30 as shown in Fig. 1. Similar results can be seen in Fig. 4. for the Ho, Er co-doped sample HE10. The crystal planes (100) and (101) in the SAD patterns in Fig. 4(b) are also consistent with the XRD pattern of HE10 indicating the ␣Sialon phase. EDS spectra 1, 2 and 3 of the points within the grains show the presence of both Ho3+ and Er3+ ions inside the grains. Conformation of the ␣-Sialon phase via XRD and SAD pattern analysis in conjunction with the result of EDS analysis of the presence of both Ho3+ and Er3+ ions inside the ␣-Sialon grains indicate that both Ho3+ and Er3+ ions have been successfully doped into the interstitial sites of the ␣-Si3 N4 matrix to form duel ions stabilized (Ho3+ , Er3+ )-␣-Sialon. The co-existence of Ho3+ and Er3+ ions into the lattice is very important for the effective energy transfer between the ions. Some Ho3+ and Er3+ ions were observed also at triple junctions (point 4) and grain boundaries (point 5) as indicated by the corresponding EDS spectra. In our previous works [31,32], all the Er3+ ions were found at grain boundaries and triple junctions while only smaller co-doped cations like Mg2+ or Yb3+ were found inside the grains.
The high-resolution transmission electron microscopy (HRTEM) image of H30 is shown in Fig. 5. A thin intergranular grain boundary phase of thickness 2.15 nm can be seen in Fig. 5(c). The calculated interplanar spacing of 0.25 nm in Fig. 5(d) corresponds to (210) plane which is the strongest ␣-Sialon peak in the XRD spectrum. The d-spacing of (101) plane in SAD pattern of Fig. 5(e) was found to be 0.44 nm identical to that in Fig. 5(b). Similarly, HRTEM image of HE10 is shown in Fig. 6. The intergranular grain boundary phase is reduced approximately to 1.50 nm in HE10, as shown in Fig. 6(c), as compared to that of the Ho singly doped sample, H30. Hence, because of co-doping of Ho3+ and Er3+ ions the amount of intergranular phase is reduced by the transient liquid phase being absorbed into the matrix of ␣-Sialon phase during sintering [5]. The interplanar spacing of 0.28 nm and 0.68 nm in Fig. 6(b) and (d) correspond respectively to the planes (002) and (100) in the XRD spectrum. We observed a very good agreement between all the calculated dspacing of the indicated lattice planes in HRTEM images of both H30 and HE10 samples with the d values of the corresponding planes of XRD spectra of JCPDS 33-261 of ␣-Sialon. Hardness, fracture toughness and density of Ho singly doped and Ho, Er co-doped samples are tabulated in Table 3. The inherent outstanding mechanical properties of ␣-Sialon ceramics are consistently retained by all the sintered samples with hardness above 20 GPa and fracture toughness above 5 MPa m1/2 . The ␣-Sialon ceramics exhibit high hardness due to their small grain size and high atomic density in unit cell structure while thin glassy phase
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Fig. 5. (a) HRTEM micrograph of H30; (b), (c) and (d) are the magnified images of the selected areas in (a); (e) is the selected area diffraction (SAD) pattern of the area (b).
Fig. 6. (a) HRTEM micrograph of HE10; (b), (c) and (d) are the magnified images of the selected areas in (a); (e) is the selected area diffraction (SAD) pattern of the area (d).
Table 3 Mechanical properties and densities of different ␣-Sialon samples.
3.2. Transmission and absorption spectra
Samples
Hardness (GPa)
Fracture toughness (MPa m1/2 )
Density (g/cm3 )
HE05 HE10 HE15 HE20 HE25 H30
20.24 20.54 20.45 20.35 20.19 20.33
5.18 5.59 5.44 5.61 5.52 5.59
3.36 3.38 3.41 3.42 3.42 3.43
that exists as intergranular phase provides high fracture toughness. The observed hardness and fracture toughness of the sintered ␣Sialon ceramics are better than that for yttrium aluminum garnet (YAG), Y2 O3 , and other optically active polycrystalline transparent ceramics available commercially [31]. Measured low densities of the samples additionally make ␣-Sialon ceramics a light weight material.
Transmission spectra of Ho singly doped and Ho, Er co-doped samples each of thickness 0.50 mm are shown in Fig. 7. The transmission measurements were performed in the wavelength range from 200 nm to 5500 nm in order to cover visible and IR regions. The transparency first increases with the increasing wavelength, reaches as high as 80% in the mid-IR region and then decreases. Zhong et al. [9] performed the first principle calculation to understand the underlying transparency mechanism in ␣-Sialon. According to them, the doped rare-earth atom interacts strongly with the doping states of ␣-Sialon, resulting in the increase in the optical gap, which suppresses the photoabsorption in the wavelength region between 1000–4500 nm and leads to the better optical transparency. The optical image of the samples in the inset shows that the sintered samples have very good transparency in visible region as well. However, Ho singly doped sample i.e., H30 has relatively poor transmittance as compared to the Ho, Er co-
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Fig. 7. Transmission spectra of all the ␣-Sialon samples of thickness 0.50 mm. Inset is the optical image of the samples.
Fig. 9. Photoluminescence spectra of different ␣-Sialon samples under 980 nm excitation. Inset (a) is the magnified view of the upconversion emission band at 766 nm, inset (b) is the down conversion spectra. The left inset is the photograph of green and red luminescence from E30 and HE05 respectively as observed by the naked eye. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
Fig. 8. Absorption spectra of Ho singly doped, Er singly doped and Ho, Er co-doped ␣-Sialon samples.
doped samples. It can be inferred that the co-doping of Ho and Er has significantly improved the optical transparency. Improvement in transparency can be attributed to the following reasons. Firstly, co-doping of Ho and Er has yielded better grain morphology as compared to the Ho single doping which is evident from Fig. 2. Secondly, the intergranular grain boundary phase, which is mostly responsible for the loss due to scattering in polycrystalline ceramics [8,10], has been reduced in the co-doped sample as is evident from Figs. 5 and 6. The reduction in the grain boundary phase obviously offers a more smooth path for light propagation across the grain boundaries. Improved optical transparency could also be due to the reduced residual pores, which act as light scattering centers, in the sintered body of the co-doped samples as compared to that in Ho3+ singly doped sample. Absorption spectra of Ho singly doped, Er singly doped and Ho, Er co-doped samples in the wavelength range of 200 nm to 2500 nm are shown in Fig. 8. All the characteristics absorption transitions of Ho3+ and Er3+ ions from their corresponding ground states 5 I8 and 4 I15/2 have been thoroughly labeled. In the absorption spectrum of H30, there is no absorption band around 980 nm. Hence,
Fig. 10. Er concentration dependence of intensities at 553, 656 and 671 nm emissions in ␣-Sialon samples.
conventional 980 nm pumping cannot be used for the Ho singly doped ␣- Sialon. However, the presence of absorption band 4 I11/2 of Er3+ at 980 nm in the HE10 indicates that the co-doped samples can be excited by 980 nm laser. Moreover, it can be seen that a pair of absorption bands, in the visible region, of Er3+ : 4 I15/2 → 4 S3/2 and 4 I15/2 → 4 F9/2 overlap that of Ho3+ :5 I8 → 5 S2 (5 F4 ) and 5 I8 → 5 F5 transitions, respectively. Hence, a sensitized upconversion luminescence [33,34] can be expected in Ho, Er co-doped ␣-Sialon. 3.3. Photoluminescence properties The room temperature photoluminescence spectra of all the samples under 980 nm excitation have been shown in Fig. 9. No upconversion emission can be seen for Ho singly doped sample, H30, because Ho3+ cannot be excited at 980 nm. For Er singly
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Fig. 11. Power dependence and the photonic process of upconversion in Ho and Er co-doped ␣-Sialon samples.
doped sample, E30, there is a strong green emission and week red and near infrared emissions. A frequency downconversion band is also observed around 1530 nm as shown in the inset (b) of Fig. 9. The emission spectra drastically change in Ho and Er codoped samples. A very intense red emission can be seen around 671 nm along with a new weak emission at 766 nm in all the co-doped samples. However, the intensities of green and NIR emissions decrease in the co-doped samples as compared to that in E30. The emissions at 535 nm, 553 nm, 671 nm, 766 nm, 812 nm and 1530 nm are assigned to Er3+ :2 H11/2 → 4 I15/2 , Er3+ :4 S3/2 → 4 I15/2 , Er3+ : 4 F9/2 → 4 I15/2 , Ho3+ : 5 S2 (5 F4 ) → 5 I7 , Er3+ :4 I9/2 → 4 I15/2 , and Er3+ :4 I13/2 → 4 I15/2 respectively. It is found that when Er concentration is increased from zero to 2.5 mol%, a new peak gradually appears at 656 nm whose intensity also increases while the intensity at 671 nm emission first increases up to 1.5 mol% and then decreases. This behavior has been depicted in Fig. 10. The bottom axis of the figure labels the Er concentration and the top axis represents the corresponding sample. We believe that at higher Er concentration energy transfer from Er3+ to Ho3+ becomes more efficient thereby increasing the intensity at 656 nm whereas the emission at 671 nm from Er3+ : 4 F9/2 → 4 I15/2 transition itself is reduced. Hence, the red emission at 656 nm can be assigned to Ho3+ : 5 F → 5 I transition. The emissions at 656 nm and 766 nm are the 5 8 strong evidence of the energy transfer from Er3+ to Ho3+ . It should be emphasized here that the green emission in E30 and red emissions in all co-doped samples are intense enough to be observed even by the naked eye at excitation power as low as 75 mW (see Fig. 9). The first reason for the observed efficient upconversion is the co-existence of Er3+ and Ho3+ ions in the lattice of ␣-Sialon. As revealed earlier in EDS analysis, Er3+ and Ho3+ ions have been successfully doped into the two large interstices of the ␣-Sialon. This makes the ions sufficiently close for effective energy transfer from Er3+ to Ho3+ . The energy transfer probability is determined by the overlapping of the wave functions of the sensitizer and activator and this probability is high only when activator and sensitizer occupy adjacent lattice sites [33]. The second reason for the efficient
Fig. 12. Simplified energy level diagram of Er3+ and Ho3+ and involved energy transfer mechanisms.
upconversion can be ascribed directly to the different crystal-fields around the trivalent lanthanide ions in matrices of ␣-Sialon. The unit cell of ␣-Sialon has hexagonal geometry and can, therefore, be considered a low symmetry host. It is known that the low symmetry hosts typically exert a crystal field containing more uneven components around the dopant ions compared to high symmetry counterparts [20]. The uneven components in ␣-Sialon matrix therefore enhance the electronic coupling between 4f energy levels and higher electronic configuration and subsequently increase f-f transition probabilities of the dopant ions giving enhanced upconversion emission. Moreover, decrease in cationic size or unit cell volume of the host can cause an increase in the crystal field strength around the dopant ions and lead to enhanced upconversion effi-
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Fig. 13. CIE chromaticity diagram of Er singly doped and Ho, Er co-doped ␣-Sialon ceramics. The top inset shows the color coordinates and the color purity of the samples.
ciency. This fact is consistent with our observation that HE05 has the strongest upconversion emission which has the least unit cell volume. The pump power dependence of luminescence intensities has been investigated to understand the upconversion mechanism. A typical behavior of HE05 for the incident power from 75 mW to 300 mW has been shown in Fig. 11(a). It has been shown that for the case of a small upconversion rates, the luminescence intensity I is proportional to the nth power of the absorbed pump power P. Mathematically [35], I ∝ P n ⇒ Log(I) ∝ nLog(P)
(1)
where n, called the order of the upconversion process, is the number of pump photons absorbed per upconverted photon and is given by the slope of the luminescence intensity versus pump power in double-logarithmic representation. The results presented in Fig. 11(b)–(f) within the above mentioned power range for all the samples show that the value of n is nearly 2. Hence, the luminescence intensity shows the quadratic dependence on the pump power indicating a two photon absorption process predominantly populates green and red emitting states. Some deviation of n from typical values can be attributed to the competition between the linear decay and the upconversion processes for the depletion of the intermediate excited states [35].
Based on the simplified energy level diagram, Fig. 12, a mechanism for the observed photoluminescence phenomena has been presented. Excited state absorption (ESA) and energy transfer (ET) upconversion are the dominant upconversion mechanisms in Er3+ , Ho3+ doped systems [28,36,37]. We first discuss the ESA mechanism of Er3+ excitation. Upon 980 nm excitation, Er3+ ions are first excited from 4 I15/2 to 4 I11/2 by the ground state absorption (GSA). On one hand, further excitation of Er3+ : 4 I11/2 takes place by excited state absorption ESA1 to make the Er3+ : 4 F7/2 level populated. The populated Er3+ : 4 F7/2 level then relaxes rapidly and non-radiatively to the next lower levels 2 H11/2 and 4 S3/2 . The final radiative transitions 2 H11/2 → 4 I15/2 and 4 S3/2 → 4 I15/2 produce green emissions at 535 nm and 553 nm respectively. On the other hand, some of Er3+ ions in 4 I11/2 nonradiatively decay to 4 I13/2 . And a subsequent radiative transition to 4 I15/2 yields 1530 nm downconversion emission. Moreover, some of the Er3+ ions in 4 I13/2 can also be excited to 4 F9/2 via ESA2 to produce either red emission at 671 nm by 4 F9/2 → 4 I15/2 transition or NIR emission at 812 nm by 4 I9/2 → 4 I15/2 after a nonradiative relaxation of the 4 F9/2 level. We now consider the energy transfer (ET) from Er3+ to Ho3+ . Here Er3+ is a sensitizer and Ho3+ is an activator. The excited Er3+ ions in 4 I11/2 state may transfer their energy to Ho3+ ions. The 5 I6 level of Ho3+ lies about 1290 cm−1 below the 4 I11/2 level of Er3+ . This energy can easily be taken up [28] by the ␣-Sialon lattice since its vibrational mode exists at 760 cm−1 . Therefore Ho3+ ions
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in the ground state 5 I8 are excited to the 5 I6 state by the neighboring Er3+ ions via ET1: 4 I11/2 Er3+ + 5 I8 Ho3+ → 4 I15/2 Er3+ + 5 I6 Ho3+ . The nonradiative relaxations of 5 I6 → 5 I7 populate 5 I7 level of the Ho3+ ions. The second step involves the excitation of Ho3+ either from 5 I6 to 5 S2 , 5 F4 state via ET2: 4 I11/2 Er3+ + 5 I6 Ho3+ → 4 I15/2 Er3+ + 5 S2 , 5 F4 Ho3+ , or from 5 I7 to 5 F5 state via a phonon-assisted ET3: 4 I11/2 Er3+ + 5 I7 Ho3+ → 4 I15/2 Er3+ + 5 F5 Ho3+ . Moreover, 5 F5 can also be populated by the nonradiative decay of 5 S2 , 5 F4 states of Ho3+ . The radiative decay of Ho3+ : 5 F5 → 5 I8 gives red emissions at 656 nm while the decay of Ho3+ : 5 S2 , 5 F4 → 5 I7 gives NIR emission at 766 nm. Because of the multiphonon relaxation, Ho3+ ions in 5 I6 are more likely to relax to lower level 5 I7 and those at 5 S2 level are more likely to relax to lower 5 F5 level nonradiatively [38]. These two processes decrease the population of 5 S2 level and increase that of 5 F5 . Therefore, the red emission is dominant and NIR emission is weak in Ho3+ . At higher Er3+ concentration, yet another energy transfer process ET4 becomes significant [24]: 4 F9/2 Er3+ + 5 I7 Ho3+ → 4 I13/2 Er3+ + 5 F5 Ho3+ . As a consequence, the red emission intensity at 656 nm due to Ho3+ : 5 F5 → 5 I8 increases while that at 671 nm due to Er3+ :4 F9/2 → 4 I15/2 decreases which were observed in the PL emission spectra of Fig. 9. Fig. 13 is the Commission Internationale de L’Eclairage (CIE) chromaticity diagram of all the samples under 980 nm excitation. The resultant CIE coordinates of the samples listed in the inset were calculated from the corresponding photoluminescence emission spectra of Fig. 9. The color purity was calculated using the method discussed in our previous work [31]. It can be seen that Er single doped sample falls in the green region while all the Ho and Er co-doped samples fall in the red region. The high color purity of the samples broadens the fields of application of this material as upconversion phosphors.
3.4. Raman spectra Raman spectra of the sintered samples are shown in Fig. 14. The sample H30 shows a broad Raman band centered around 650 cm−1 . The inset shows the Gaussian resolved components of the sample H30 with the highest vibrational band at 800 cm−1 . The fundamental structural units in Si3 N4 crystal and Sialons are the Si–N4 tetrahedra, linked in such a way that the N– Si3 triangular unit is approximately planar. Therefore, the vibrational spectra of Sialons usually display a broad band attesting to the significantly distorted structure of these Si– N4 tetrahedra and N–Si3 triangular units [39]. Spectra of Ho, Er co-doped samples are significantly different from that of Ho single doped sample. It is evident that the incorporation of Er3+ ions splits the peaks into two distinct low and high frequency bands below and above 650 cm−1 . First principle calculations [40–42] of vibrational density of states of ␣-Si3 N4 have also shown that two different modes of vibration of N atoms give rise to the well separated bands below and above 600 cm−1 . Si N bond is much stiffer against stretching than bending [42]. Low and high frequency phonon bands below and above 650 cm−1 in our case can, therefore, be ascribed respectively to bond bending (rocking motion) and bond stretching vibrations of N atoms in N– Si3 triangular units. The very low frequency band at 135 cm−1 is attributed to the vibrations of the acoustic and optic branches distorting the whole Si N bonding network. The observed vibrational frequencies of 468, 515, 545, 760 and 855 cm−1 are in good agreement with other reports [43,44] of experimental and calculational results for ␣-Si3 N4 . The highest phonon energy of 855 cm−1 for Ho, Er co-doped ␣Sialon is still less than that of the well known upconversion and laser host materials such as phosphate glass (1200 cm−1 ), Silica glass (1100 cm−1 ), LaPO4 (1050 cm−1 ), YAG (860 cm−1 ), and fluorophosphate glass (1060 cm−1 ) [18,26]. The moderately low value
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Fig. 14. Raman Spectra of different ␣-Sialon samples. The inset shows the Gaussian resolved components of Raman spectrum of H30 sample.
of phonon energy has the effect of reducing the non-radiative multiphonon relaxation in the ␣-Sialon host and enhancing the novel energy transfer process between Er3+ and Ho3+ ions for the efficient photoluminescence. 4. Conclusions We have demonstrated a very intense infrared to visible frequency upconversion and near infrared frequency downconversion photoluminescence for the first time in Ho3+ and Er3+ co-doped transparent ␣-Sialon ceramics, sintered by hot press technique, under 980 nm laser excitation. The intense upconversion emission was due to the efficient energy transfer process between Er3+ and Ho3+ ions as a result of co-existence of the ions in the interstices of the ␣-Sialon lattice. Formation of the ␣-Sialon phase and existence of both Er3+ and Ho3+ ions in the lattice were confirmed by XRD, TEM, and EDS analysis. Transparency of the sintered samples was improved by the co-doping of Er3+ and Ho3+ ions due to the reduced intergranular boundary phase in the co-doped samples. Raman spectroscopic study revealed that the ␣-Sialon ceramics have moderately low phonon energy as compared to some well known phosphate and oxide glasses and ceramics. The inherent outstanding thermo-chemical and mechanical stability along with the up and downconversion emission properties and moderately low phonon energy of (Ho3+ , Er3+ )-␣-Sialon ceramics make the material promising candidate for multifunctional applications. Acknowledgment This research was supported by Global Research Laboratory (GRL) program of the National Research Foundation of Korea (NRF) funded by the Ministry of Education, Science and Technology (MEST), Republic of Korea grant number (2010-00339). References [1] G.Z. Cao, R. Metselaar, ␣’-Sialon ceramics: a review, Chem. Mater. 3 (1991) 242–252. [2] S. Hampshire, H.K. Park, D.P. Thompson, K.H. Jack, ␣’-Sialon ceramics, Nature 274 (1978) 880–882. [3] C. Santos, K. Strecker, P.A. Suzuki, S. Kycia, O.M.M. Silva, C.R.M. Silva, Stabilization of ␣-SiAlONs using a rare-earth mixed oxide (RE2 O3 ) as sintering additive, Mater. Res. Bull. 40 (2005) 1094–1103.
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Please cite this article in press as: Y.K. Kshetri, et al., Intense visible upconversion emission in transparent (Ho3+ , Er3+ )-␣-Sialon ceramics under 980 nm laser excitation, J Eur Ceram Soc (2016), http://dx.doi.org/10.1016/j.jeurceramsoc.2016.07.005