Nano-glass–ceramics containing chromium-doped LiGaSiO4 crystalline phases

Optical Materials 32 (2010) 896–902

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Nano-glass–ceramics containing chromium-doped LiGaSiO4 crystalline phases K.A. Subbotin a,*, V.A. Smirnov a, E.V. Zharikov b, L.D. Iskhakova c, V.G. Senin d, V.V. Voronov a, I.A. Shcherbakov a a

A.M. Prokhorov General Physics Institute of the Russian Academy of Sciences, Vavilova St. 38, Moscow 119991, Russia D.I. Mendeleyev University of Chemical Technology of Russia, Miusskaya Sq. 9, Moscow 125047, Russia c Fiber Optics Research Centre of the Russian Academy of Sciences, Vavilova St. 38-F, Moscow 119991, Russia d V.I. Vernadsky Institute of Geochemistry and Analytical Chemistry of the Russian Academy of Sciences, Kosygina St.19, Moscow119991, Russia b

a r t i c l e

i n f o

Article history: Received 19 August 2009 Received in revised form 24 December 2009 Accepted 9 January 2010

Keywords: Nano-glass–ceramics Chromium (IV) Spectroscopy

a b s t r a c t Chromium-doped nano-glass–ceramic samples containing crystallites of a- and c-polymorphous modifications of eucryptite-like LiGaSiO4 were fabricated by devitrification of Cr–Li–Ga–Si–O glass of various compositions. Evaluated sizes of the crystallites in synthesized samples are from tens of nanometers (in the case of ceramming at 570–650 °C during 2 h) up to tens of microns (additional heat treatment at 900 °C during 5 min). Spectroscopic investigations (absorption and fluorescence spectra, fluorescence decay kinetics) of the prepared vitreous precursors and of the nano-glass–ceramic samples were performed. The absorption bands of Cr3+ and Cr6+ ions were observed in the vitreous precursors, whereas the absorption of Cr4+ ions prevails in the samples after ceramming. The measured fluorescent properties of the glass–ceramics containing only a-LiGaSiO4:Cr4+ crystallites are as follows: the fluorescence peak position is 1280 nm, the FWHM of the fluorescence band is 270 nm, the lifetime of the excited 3T2 (Cr4+) state is 12 ls at 300 K and 65–70 ls at 77 K. These properties are similar to those of the earlier investigated Cr:LiGaSiO4 single crystals. One more Cr4+ fluorescent center with a shorter decay time (1– 1.5 ls at room temperature and 7–9 ls at 77 K) was observed in the samples in which c-LiGaSiO4:Cr4+ crystallites were present. The origin of this fluorescence is so far unclear. Ó 2010 Elsevier B.V. All rights reserved.

1. Introduction Since the first demonstration of laser action of tetrahedrallycoordinated tetravalent chromium in forsterite Mg2SiO4 [1,2] and in yttrium–aluminum garnet (YAG) Y3Al5O12 [3] at the end of the 80-ies of the XX century, Cr4+-based tunable solid-state lasers attract significant attention from the point of view of both basic investigations and practical applications [4,5]. However, up to now only Cr4+:YAG and Cr4+:forsterite crystals have found commercial application as Cr4+-laser media. Unfortunately, these crystals have a rather low fluorescence quantum yield (22% for Cr4+:YAG and 9–16% for Cr4+:forsterite) and a relatively short lifetime of the 3T2 upper laser level of Cr4+ (4 and 2.7 ls, respectively) [2,6,7]. Besides, it is impossible to achieve significant Cr4+ concentrations in both these materials. Search for alternative, more efficient laser hosts for Cr4+ has been performed by numerous scientific groups during the past 20 years. This activity has resulted in the development of various promising crystals (see, for example, review [4]). However, neither of these crystals can be grown directly from melt as bulk samples * Corresponding author. Tel.: +7 499 503 83 95; fax: +7 499 135 02 70. E-mail address: [email protected] (K.A. Subbotin). 0925-3467/$ - see front matter Ó 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.optmat.2010.01.016

with laser quality because of the following factors: phase transitions below the melting point, and/or incongruent melting of the compound, and/or high selective evaporation of separate melt components, etc. One of the above-mentioned alternative promising crystals is aeucryptite-like Cr:LiGaSiO4 discovered and studied by our team [8]. The estimated value of the fluorescence quantum yield of Cr4+ in this crystal is 28% [9], whereas the measured fluorescence lifetime is 14 ls at 300 K. Moreover, owing to the absence of octahedral cation sites in the structure of this crystal [10], the formation of essential amounts of parasitic divalent and trivalent chromium ions must not occur here. However, incongruent melting of the compound [11–13] prevents the growth of high-quality Cr:LiGaSiO4 bulk single crystals directly from melt. Phase transformations in some areas of the ternary system Li2O–Ga2O3–SiO2, including the composition Li2O:Ga2O3:SiO2 = 1:1:2 (hereafter we used molar ratios) were studied in paper [13]. The possibility to prepare LiGaSiO4 glass–ceramics was shown in that paper for the first time. It was found that heat treatment of glass with composition Li2O:Ga2O3:SiO2 = 1:1:2 at temperatures of 600–650 °C results in the formation of metastable c-LiGaSiO4 crystalline phase. This phase irreversibly transforms into stable eucryptite-like a-LiGaSiO4 phase during additional heat treatment

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at 900 °C for 5 min. The transformation also occurs at lower temperatures, but with substantially lower rates. Transparent nano-glass–ceramics has advantages of both crystalline and vitreous laser materials and is an appropriate alternative to single-crystalline laser media. Glass–ceramics can have almost the same spectroscopic properties as the corresponding single crystals. On the other hand, nano-glass–ceramics can be fabricated using relatively simple techniques, as compared to many single crystals. Nano-glass–ceramics can be fabricated as a massive sample, or as a single-mode fiber waveguide. The latter is very important for the possibility to use these materials in lasers and amplifiers in telecommunication [14,15]. Up to now, there have been only a few papers devoted to the investigation of Cr4+-doped transparent nano-glass–ceramic materials as potential laser gain media. In particular, transparent glass– ceramics containing chromium-doped forsterite nano-crystals was fabricated and investigated in Refs. [15,16]. The spectroscopic properties of this material are comparable with those of bulk Crforsterite crystals. Transparent Cr4+-doped gelhenite nano-glass– ceramics (Ca2Al2SiO7) was studied in Ref. [17]. Glass–ceramics containing chromium-doped willemite Zn2SiO4 crystallites was investigated in Ref. [18]. Note that willemite possesses the structure related to that of a-LiGaSiO4. However, the Cr4+ fluorescence decay times in these materials are, again, rather short (only 1–3 ls, close to that for the corresponding single crystals). Thus, these glass– ceramic materials are not very interesting for practical applications. Glass–ceramics containing eucryptite LiAlSiO4 crystallites (the structure is related to that of LiGaSiO4) is well-known. It has found practical applications in various optical and in other devices, in which a low thermal expansion coefficient is the key parameter. LiAlSiO4 nano-glass–ceramics doped by some transition and rareearth ions (e.g. Nd3+ and Co2+) is also being studied as active and passive laser media [19,20]. However, in the available literature we failed to find evidence on successful fabrication of transparent LiGaSiO4 nano-glass–ceramics and on its evaluation as a potential laser medium performed by other teams. We are carrying out the investigations on the synthesis of Cr4+:LiGaSiO4-based nano-glass–ceramics, as well as its optical and spectroscopic characterization during the last few years [21– 23]. In this paper, we summarize some our results of these investigations.

2. Experimental The experimental samples were fabricated by a rather conventional glass–ceramics technology, which includes the following stages:  Fabrication of the parent glass by melting the charge and by quenching the melt;  Two-step heat treatment of the fabricated glass. During the first step, heating is performed at a lower temperature, and nucleation occurs at this step. Then, an increase of temperature leads to growth of the nuclei up to the required sizes. For the preparation of the charges, the initial chemicals – Li2CO3, Ga2O3, SiO2 and Cr2O3 of extra-pure grade - were weighted in appropriate ratios on the electronic abacus ‘Sartorius’ with the precision of ±0.001 g. Thereafter, 0.2 wt.% of Cr2O3 was added to all charges. This amount of Cr2O3 corresponds to the concentration of Cr ions in the melt of 5.4  1019 cm3. The weighted chemicals were thoroughly mixed. Then the powders were pressed into cylindrical feed-rods with a diameter of 8 mm and a length of 12 cm.

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The feed-rods were sintered during 24 h at a temperature of 980 °C. Melting of the feed-rods was performed in the ‘URN-2-ZP’ furnace with light heating. Normally, this furnace is used for the growth of refractory single crystals by the container-free floating zone melting technique. Details of the ‘URN-2-ZP’ furnace design, as well as an information on the growth features for a number of technically important crystals are given in Ref. [24]. Our experience on the Cr:LiGaSiO4 single crystal growth with the help of this furnace is described in Ref. [11]. In this study, we melted the lower tip of the feed-rod, until a transparent and uniform drop of melt was formed. Because the Cr:LiGaSiO4 melt is usually transparent, we were able to control visually the uniformity of the melt drop and completeness of its melting. The drop was hanging at the tip of the feed-rod due to surface tension. A crucible was not used in this method, and hence, no problems occurred with contamination of the melt by the material of the crucible. The quenching of the melt drop was performed by two different methods: (i) by abrupt termination of the heating light delivery to the melting zone. The heat removal from the drop occurs in this case only by heat radiation and by air convection, and the evaluated average rate of temperature reduction is 100– 200 °C per second at the very beginning of the quenching; (ii) by dropping the melt onto a massive copper bar. In this case, the main way of heat removal from the drop is heat conductivity towards the copper bar. The evaluated average rate of temperature reduction increases in this case up to 400– 500 °C per second at the very beginning of the quenching. The vitreous precursors were devitrified by heat treatment in an isothermal resistive furnace. Temperature control was performed by a chromel–alumel thermocouple with the precision of ±1 °C. The duration of the nucleation step was from 24 to 120 h for different samples, the duration of the growth stage being 2 h. The microprobe analysis of the vitreous and glass–ceramic samples was carried out with the help of ‘Cameca SX-100’ analyzer, in order to measure the actual chromium concentration in the samples. The concentration for each sample was calculated as the average value of five probes at five different points of the sample slice. The analyzer was coupled to ‘JSM-5910LV’ scanning electron microscope (SEM), which allowed us to observe directly the number and the sizes of crystallites in the samples. In order to obtain an appropriate SEM contrast between the vitreous and crystalline phases (which possess almost the same chemical composition), we etched polished slices of the samples in aqueous solution of HF + NH4HF2 mixture. The rate of etching for glass is substantially higher than for crystallites. This fact leads to the appearance of an acceptable topological contrast in the SEM photographs when applying the scanning regime of back-scattered electrons. The average sizes of crystallites in cerammed samples were also evaluated from the X-ray powder diffraction (XRD) pattern with the use of the Scherrer equation [25]:

D¼

kk ; cos h  b

ð1Þ

where k is the wavelength of the probe X-rays; h is the peak position of XRD reflex; b is the width (FWHM) of this reflex, and k is the correction factor, which takes into account the thermal stresses at the boundaries of crystallites and inside them. Note that heat treatment of our samples takes a rather long time (see above) with rather low heating and cooling rates (several tens of degrees per hour). We believe, these conditions were sufficient for significant relaxation of

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the stresses to a large extent. Thus, in evaluative calculations, we did not take into account the impact of thermal stresses and used k = 1. XRD measurements of our samples were carried out on a DRON4 diffractometer using Cu Ka radiation, k = 1.54 Å. For calculations by Eq. (1), we used only strong enough XRD reflexes, which have a sufficiently good resolution on the background of the diffuse halo of the residual vitreous phase. To measure the optical absorption spectra, the samples were prepared in the form of platelets with a thickness varying from 0.5 to 2 mm. The platelets were polished on both sides. The measurements were made on a ‘Shimadzu UV-3101PC’ spectrophotometer in the spectral range 300–2000 nm with a resolution of 1 nm. During the calculations of the absorption coefficients, the reflection from polished facets of the platelets were taken into account. The spectral and kinetics properties of the Cr4+ fluorescence in the samples were studied upon excitation by a Q-switched Nd:YAG laser with a pulse duration of 100 ns. The fluorescence signal was dispersed by an ‘MDR-2’ monochromator (the dispersing element was a grating with 300 grooves per mm), detected by a photomultiplier and then the electrical signal was amplified and passed onto a digital oscilloscope. The fluorescence spectra were recorded with a resolution of 1 nm. The lifetimes were calculated from the decay kinetics, which was approximated in the framework of the twoexponential decay model by means of the MicrocalÒOriginÒ 6.0. software. The accuracy of the lifetimes calculations was better than ±0.5 ls.

3. Results and discussion 3.1. Synthesis and characterization of vitreous precursors We fabricated compositions:

vitreous

precursors

of

three

different

(a) stoichiometric composition, Li2O:Ga2O3:SiO2 = 1:1:2; (b) composition with slight deficiency of silica, the main glass– forming component in this system, Li2O:Ga2O3:SiO2 = 1:1:1.95; (c) composition with slight excess of silica, Li2O:Ga2O3:SiO2 = 1:1:2.1. During charge melting on the ‘URN-2-ZP’ setup, transparent homogeneous melts were usually formed within 2–3 min after starting the feed-rod heating. In our first experiments, the melts were quenched immediately after finishing their formation (i.e. after complete dissolution of all visible inclusions).

The vitreous precursors quenched by method (ii) usually contained knags, cracks, and other bulk defects. Thus, this quenching method is far short of optimum. However, in our first experiments with composition (b), only method (ii) was the only available way of vitrification the melt, without its partial spontaneous crystallization. Fortunately, it was found later that keeping a melt above the melting point during approximately 30 min drastically simplifies vitrification. This fact can be explained by gradual polymerization of [SiO4] tetrahedra, a phenomenon known for silicate melts [26] and promoting their vitrification. After such prolonged melting, we succeeded in obtaining high-quality vitreous samples of all compositions studied via method of quenching (i). The samples vitrified after such prolonged melting, possessed a sufficiently good optical quality with a yellowish-green color, and did not contain cracks, knags, visible (or detectable by XRD analysis) inclusions of crystalline phase, gas bubbles, etc. (Fig. 1a). Prolonged heating, however, leads to increased chromium evaporation losses from the melt. Microprobe analysis of two samples of vitreous precursors, which were obtained from the same feedrod with different durations of melting, showed that the total chromium content for sample #1 (long melting) is (0.80 ± 0.04)  1019 cm3, which is by a factor of 3 less than that for sample #2 ((2.3 ± 0.1)  1019 cm3, short melting), and more than by a factor of 6.5 less than that in the initial charge (see above). The optical absorption spectra of these two samples are shown in Fig. 2. Two oxidation states of chromium can be easily identified in both the samples from these spectra: tetrahedrally-coordinated Cr6+ (the absorption band at 350 nm which can be attributed to the ligand-to-metal charge transition [27]), and octahedrally coordinated Cr3+ (the absorption peak at 630 nm which can be attributed to the 4A2 ? 4T2 transition, and a shoulder at 450 nm arising owning to the 4A2 ? 4T1 transition). These oxidation states are typical for chromium in silicate glasses [27,28]. It is seen from Fig. 2 that a decrease of Cr6+ and Cr3+ ions concentrations in Cr:LiGaSiO4 glass with increasing the melting duration is not the same: the Cr6+ content is almost five times smaller in sample #1, whereas the corresponding reduction of the Cr3+ content is less than twofold. This distinction in the behavior of Cr6+ and Cr3+ ions in the Li– Ga–Si–O melt can be explained by the following reasons:  Higher volatility of Cr6+ in comparison with Cr3+. The dependencies of the chromium-containing vapor pressures above the CrO3 and Cr2O3 oxides were studied in Refs. [29] and [30], respectively. One can conclude from these dependencies that the vapor pres-

Fig. 1. Example of fabricated chromium-doped Li–Ga–Si–O glass (a), and of glass–ceramics (b).

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Fig. 2. The absorption spectra of two vitreous precursors obtained from the same feed-rod of composition (c) with different durations of melting: #1, 30 min; #2, 3 min.

sures above CrO3 and Cr2O3 are the same (0.04 Pa) at temperatures of 150 °C and 1550 °C, respectively. Extrapolation of these dependencies to the temperature of LiGaSiO4 melt (1300 °C) gives a difference in vapor pressures above CrO3 and Cr2O3 of 17 orders of magnitude! Of course, correctness of such an extrapolation is rather questionable. Nevertheless, a drastic difference in volatilities of Cr6+ and Cr3+ is evident from the discussed data.  Gradual Cr6+ ? Cr3+ transformation at high temperatures. It is well-known (see, for example, [31]) that for ions with alternating oxidation states, thermodynamic stability of more reduced states in the same ambient increases with increasing temperature, whereas stability of more oxidized states, on the contrary, diminishes. The fields of thermodynamic stability of different chromium oxides are given in Ref. [32] and fully confirm the above tendency. 3.2. Devitrification of precursors The data from the literature [12,13], our DTA analysis, and the preliminary ceramming experiments allowed us to determine the optimal temperatures of the nucleation (550–570 °C) and crystallization (600–610 °C) stages of devitrification of the glass investigated. These optimum temperatures are almost the same for all compositions studied. However, the rates of nucleation and crystallization evaluated from the SEM and XRD results decrease with increasing the SiO2 fraction in the glass composition (‘‘b” ? ‘‘a” ? ‘‘c”): the nucleation rate decreases by an order of magnitude, and the crystallization rate, by several tens of percent. In our devitrification experiments we tested a little wider temperature ranges of the ceramming process in comparison with above optimal ones, i.e. 530–600 °C for the nucleation and 570–650 °C for the growth stage. After devitrification, all the samples became semitransparent or translucent, with more or less value of turbidity (see Fig. 2b). We performed XRD analysis of the samples. In the samples underwent the lower crystallization temperatures (570–630 °C), all the XRD peaks observed can be attributed to the metastable c-LiGaSiO4 phase (JSPDS card # 39–0279). The formation of just this crystalline phase could be expected, according to the data of Ref. [13]. Traces of stable a-LiGaSiO4 phase became detectable by XRD analysis only starting form the growth stage temperature of about 640 °C From the XRD pattern we evaluated the sizes of the crystallites formed using Eq. (1). The calculation results for two samples are

given in Table 1. The deviation of the crystallites sizes calculated from different peaks of the same pattern is several tens of percent. Such a deviation is acceptable for estimative calculations. Crystallites in sample #3 have a larger average size than those of sample #4. This can be explained by a reduction of the SiO2 content in the precursor, from which sample #3 was cerammed, in comparison with the precursor used for sample #4 (see Table 1). Such changes in composition must result in an increase in the crystallization rate. Besides, it should be noted that sample #3, unlike sample #4, has much fewer number of distinct XRD peaks than sample #4, and these peaks are weaker. This fact, apparently, is due to a smaller share of the crystalline phase (lesser extent of devitrification) in sample #3. Thus, sample #3 contains a rather small number of large crystallites, whereas sample #4, much more. The nucleation temperature of sample #3 (608 °C) is outside the specified optimal range, and, therefore, the nucleation rate for this sample is, in our opinion, smaller. Therefore, despite much higher nucleation rates for SiO2-deficient glasses in comparison with those for SiO2-rich ones at the same temperature, the shift from the optimum temperature appears to be a stronger factor in this case. The crystallites sizes evaluated from the XRD data are confirmed by the SEM images, (e.g. see Fig. 3). Note that the crystallites visible at the SEM images are slightly etched. That is why the seeming sizes of the crystallites are somewhat less than actual sizes of the crystallites situated in bulk of the samples. For the same reason, we cannot see faceting of the crystallites in Fig. 3. As it was noted above, all the fabricated glass–ceramic samples have pronounced turbidity. Hence, the average sizes of the crystallites formed in the samples were too large. The data in the literature on the admissible sizes of crystallites providing acceptable turbidity of glass–ceramics are rather contradictory. In Ref. [15], the sizes of forsterite crystallites in K–Mg–Al– Si–O parent glass were up to 40 nm (i.e. comparable with our results). The nano-glass–ceramics in Ref. [15] had the total loss (absorption + scattering) in the spectral range of the Cr4+ emission (1200–1700 nm) well below 0.05 dB/cm (0.11 cm1). However, according to the absorption spectrum given in Ref. [15], the scattering loss of the same samples at shorter wavelengths, e.g. near 550 nm, is much greater, 2.5–3 cm1. Paper [14] specifies more rigid requirements for the crystallite sizes in oxyfluoride nano-glass–ceramics: < 15 nm. However, such nano-glass–ceramics has much less scattering losses: well below 100 dB/km in the range of wavelengths of 1300–1600 nm. Thus, the threshold size of crystallites is, apparently, different for different nano-glass–ceramics.

Table 1 Evaluation of grain size of two glass–ceramic samples by Scherrer equation (k = 1.54 Å). Reflex peak position, 2h (°)

Reflex half-width, b (rad.)

Calculated average size d (nm)

Sample #3 (composition ‘‘b” (Li2O:Ga2O3:SiO2 = 1:1:1.95), method of quenching (i)) 20.8 0.0023 69 22.0 0.0037 42 22.4 0.0035 45 Average

52

Sample #4 (composition ‘‘c” (Li2O:Ga2O3:SiO2 = 1:1:2.1), method of quenching (i)) 20.8 0.0048 33 22.1 0.0039 40 22.4 0.0038 41 35.2 0.0063 26 36.0 0.0054 30 Average

34

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Fig. 3. A SEM photograph of the slice of sample #4. Fig. 4. The absorption spectra before (dotted line) and after (solid line) ceramming.

In Ref. [14], some additional requirements for the crystallites are also formulated, in particular: – compatibility of the crystallite sizes with the inter-particle spacing; – narrow particle-size distribution; – absence of clustering of crystallites. One can see from Fig. 3 that our LiGaSiO4 glass–ceramics meets these requirements only partially. Thus, the ceramming regimes of our Cr:LiGaSiO4 nano-glass–ceramics require further optimization, in order to lower the turbidity. Finally, the turbidity of glass–ceramics depends very strongly on the refraction indices difference between the crystallites and the parent glass. According to [33], turbidity is proportional to (ncryst  nglass)2. Unfortunately, the corresponding data on either LiGaSiO4 crystals, or Li–Ga–Si–O glass of the compositions used in this study are unavailable. Thus, we cannot evaluate the expected turbidity in our samples in this way. However, it is clear that modification of the precursor composition with the aim to bring the refractive index of the precursor closer to that of the Cr:LiGaSiO4 crystallites would reduce the turbidity of the glass– ceramics.

3.3. Spectroscopic investigation of cerammed samples Besides the appearance of pronounced turbidity, the color of the samples after ceramming became deep blue (Fig. 1b)1. The reason for that phenomenon is drastic change in the absorption spectrum (see Fig. 4). Absorption of tetrahedrally-coordinated Cr4+ ion arises: a strong, orbitally split band between 570 and 750 nm, which can be attributed to the 3A2 ? 3T1 vibronic transition [34], and a less intense band in the range 1000–1400 nm, which can be attributed to the symmetry-forbidden 3A2 ? 3T2 vibronic transition [35]. Simultaneously, the intensities of the above bands of Cr6+ and 3+ Cr ions considerably decrease in the cerammed samples. From this fact we can draw a preliminary conclusion that Cr4+ arises in the crystallites owing to the interaction between Cr6+ and Cr3+ ions. Unfortunately, a detailed mechanism of Cr4+ formation from other chromium oxidation states during crystallization of either melts or glasses is not virtually discussed in the literature. 1 For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.

Fig. 5. The absorption spectra of a nano-glass–ceramic sample studied in this paper and a Cr:LiGaSiO4 single crystal studied earlier.

The absorption spectrum of the synthesized nano-glass–ceramic samples in comparison with the spectrum of the Cr:LiGaSiO4 single crystal studied earlier [8] are given in Fig. 5. In general, the view of these spectra is rather similar: in both cases strong Cr4+ bands are dominating. However, there are substantial differences between these absorption spectra in both peak positions, and FWHM of the Cr4+ absorption bands. These differences are, apparently, due to the fact that Cr4+ ions in our nano-glass–ceramics are located mainly in the metastable c-LiGaSiO4 phase, whereas single crystals consist of Cr4+-doped stable a-eucryptite-like phase. Unfortunately, it is not quite clear how different are the structures of a- and c-LiGaSiO4, because the latter has been investigated very poorly. However, we may conclude that the local environment of Cr4+ is substantially different in these two structures. This circumstance leads to differences in strength and symmetry of the crystal field at the Cr4+ site, and, hence, to the observed differences in the absorption spectra. The Cr4+ fluorescence in the glass–ceramic samples studied is rather weak at 300 K. That is why we failed to measure the fluorescence spectrum of the samples at 300 K with a satisfactory accuracy. The decay kinetics for most cerammed samples is close to a

K.A. Subbotin et al. / Optical Materials 32 (2010) 896–902

single-exponential law with the lifetime of 1–1.5 ls, although in some cases the contribution of a longer-lived fluorescence is also considerable. The decay time constant of the longer-lived fluorescence is 11.5–12 ls. This is close to the fluorescence decay time of a-Cr4+:LiGaSiO4 single crystal at 300 K (14 ls). At 77 K the fluorescence is much stronger, than at 300 K. The fluorescence spectrum of one such cerammed sample at 77 K is shown in Fig. 6 along with the fluorescence spectrum for a Cr4+:LiGaSiO4 single crystal at the same temperature. These spectra are rather similar, although slight discrepancies in both the peak positions and the FWHM are evident. However, again, a very large difference was observed in the fluorescence decay kinetics at 77 K. For most cerammed samples studied, the fluorescence decay kinetics at 77 K is two-exponential with a shorter lifetime of 7–9 ls and a longer lifetime of about 65–70 ls. The longer lifetime is close to the fluorescence lifetime of a Cr4+:LiGaSiO4 single crystal at 77 K (83 ls [8]). Thus, the observed longer-lived contribution to the fluorescence of the glass–ceramics studied is, apparently, due to slight amounts of the a-eucryptite phase of Cr4+:LiGaSiO4, which was observed by the XRD in some of the cerammed samples. The emitting center responsible for this fluorescence has, apparently, a structure similar to that in a-LiGaSiO4:Cr single crystal, i.e. a Cr4+ ion located in one of Si4+ or Ga3+ tetrahedral positions. The structure of the Cr4+ emitting center with a shorter lifetime remains unclear. Maybe, this center is Cr4+ located in the metastable c-LiGaSiO4 crystalline phase. Further investigation is required to clarify this point. Anyway, this emitting center is not very interesting from the point of view of practical application in solid-state lasers, in contrast to the longer-lived Cr4+ emitting center. Therefore, it is necessary to somehow increase the share of longer-lived Cr4+ emitting centers and to diminish the share of shorter-lived centers. As was noted above, metastable c-LiGaSiO4 irreversibly transforms very quickly into the stable a-eucryptite-like phase during heating at 900 °C [13]. We performed such additional short-term heating of our samples. In fact, according to the subsequent XRD analysis of the samples, the c-phase had completely disappeared, while the XRD reflexes of a- LiGaSiO4 had become much stronger.

901

Fig. 7. The normalized fluorescence spectra of an additionally heated nano-glass– ceramic sample (thin blue line) and a Cr4+:LiGaSiO4 single crystal investigated in Ref. [8] (thick red line) at room temperature. In both cases, a Q-switched Nd-YAG laser was the excitation source.

The fluorescence intensity of these samples had also drastically increased. The shape of the Cr4+ fluorescence spectra of the additionally heated glass–ceramics is quite similar with that of a single crystal at 300 K (Fig. 7). The fluorescence lifetime of these samples became 12 ls with nearly a single-exponential decay, i.e. longerlived Cr4+ emitting center became the major one in the samples after additional heating. All these facts confirm that the observed longer-lived fluorescent center is Cr4+ ion located in crystallites of the stable a-LiGaSiO4 phase. Note that this decay time is slightly less than that for Cr4+doped a-LiGaSiO4 single crystal [8]. A similar reduction of the Cr4+-fluorescence lifetime of glass–ceramics in comparison with that of single crystals was also found for Cr-forsterite at 300 K [15]. The authors of [15] associate this reduction with either the presence of two non-equivalent Cr4+ luminescent centres, or nonradiative energy transfer from Cr4+-ions in nano-crystals towards the glass host by an inductive-resonant mechanism. Thus, nano-glass–ceramics fabricated in this work containing Cr4+-doped a-LiGaSiO4 crystallites reproduce the fluorescent properties of the Cr4+-doped a-LiGaSiO4 single crystal. Unfortunately, an additional heat treatment of the samples at 900 °C during 5 min results in complete loss of their transparency because of a fast uncontrollable increase of the crystallites sizes. These crystallites become visible even with an optical microscope, their sizes being tens of microns. Therefore, the main task to be solved at the next stage of our investigations is the search for the ceramming conditions, which provide the formation of a greater number of nuclei. In this way the undesirable growth of crystallites can be prevented. Another option could be the search for techniques of faster phase transformation from c- to a-eucryptite-like phase of Cr4+:LiGaSiO4, which are not accompanied by a loss of transparency. These techniques could be as follows: – use of lower temperatures with a longer duration of the process; – use of glasses with different compositions, including the addition of modifying co-dopants (Ti4+, Zr4+, Ba2+, etc.) into the initial glass, which could promote nucleation and crystallization of aeucryptite-like phase directly from the glass phase.

Acknowledgements Fig. 6. The normalized fluorescence spectra of a nano-glass–ceramic sample studied in this paper and a Cr4+:LiGaSiO4 single crystal investigated in Ref. [8] at 77 K. In both cases, a Q-switched Nd-YAG laser was the excitation source.

The authors are grateful to Dr. D.A. Nikolayev for his help in optical absorption measurements. This work was supported by

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