Tunable upconversion emission in NaLuF4–glass-ceramic fibers doped with Er3+ and Yb3+

Novel glass-ceramic optical fibers containing NaLuF4 nanocrystals doped with 0.5ErF3 and 2YbF3 (mol%) have been prepared by the rod-in-tube method and controlled crystallization. NaLuF4 nanocrystals with a size around 20 nm are obtained after heat treatment at 600 °C. Intense upconverted green and red emissions due to (2H11/2, 4S3/2) → 4I15/2 and 4F9/2 → 4I15/2 transitions, respectively, together with a blue emission due to 2H9/2 → 4I15/2 transition have been observed under excitation at 980 nm. The intensity of the green and red upconversion bands shows a nearly linear dependence on the excitation power which can be explained by saturation effects in the intermediate energy states and proves that a sensitized energy transfer upconversion process is responsible for the population of the emitting levels of Er3+ ions. The upconversion emission color changes from yellow to green by increasing the excitation power density which allows to manipulate the color output of the Er3+ emission in the glass-ceramic fibers. The tunable emission color is easily detected with the naked eye. This interesting characteristic makes these glass-ceramic fibers promising materials for photonic applications.


Introduction
Oxyuoride glass-ceramics (OxGCs) have received great attention since the preparation of transparent materials in the 1990s. 1 During the last 7 years, several review papers regarding these materials were published. [2][3][4][5][6][7] Higher photoluminescence (PL) efficiencies and better energy-transfer (ET) are usually obtained in OxGCs with respect to glasses due to the controlled crystallization of low phonon-energy uoride nano-crystals (NCs) that are good hosts for Rare-Earth (RE) ions. [8][9][10][11][12][13][14][15][16][17][18] Most glass systems are based on aluminosilicate compositions that show good mechanical, thermal and chemical properties together with improved PL features of uoride NCs doped with RE ions. Most studies were performed on bulk samples but the possibility to prepare GC bers is also becoming a hot-spot for novel optical materials.
The rst studies about transparent GC bers date back to the end of the 1990s when P. A. Tick indicated four conditions to achieve this goal: 19 (1) crystal size smaller than 15 nm; (2) interparticle spacing comparable with the crystal size; (3) narrow particle size distribution and (4) absence of clustering. Some years later, Tick et al. produced bers with a core of 5 mm using the double crucible method starting from a Wang and Ohwaki's modied composition in the system SiO 2 -Al 2 O 3 -PbF 2 -CdF 2 -YF 3 . 20 Aer the crystallization process at 440-470 C, a crystalline YF 3 -CdF 2 -PbF 2 solid solution with an average size $6.5 nm was observed in the GC ber. The authors concluded that low losses, $1 dB m À1 or less, are possible in such systems in spite of a relatively large refractive index mismatch (0.06). In 2001, Samson et al. 21 described the rst efficient GC ber laser doped with Nd 3+ using the same glass system and drawing process reported by P. A. Tick. 20 Similar slope efficiencies $30% were obtained for glass and GC bers, but the presence of NCs embedded within the core of the singlemode optical ber allowed enhancing the uorescence and gain spectrum. R. Lisiecki et al. 22 and Augustyn et al. 23 prepared GC bers starting from the glass composition (mol%) 48SiO 2 -11Al 2 O 3 -7Na 2 O-10CaO-10PbO-11PbF 2 doped with 3ErF 3 and co-doped with 1ErF 3 and 3YbF 3 . The bers were drawn at 720-740 C using the rod-in-tube method, and then crystallized at 700 C. Pb 5 Al 3 F 19 , Er 4 F 2 O 11 Si 3 and Er 3 FO 10 Si 3 crystalline phases were obtained with size below 10 nm. The incorporation of RE dopants in the crystal phases was indicated by a narrowing of the emission spectra and longer lifetime of the rst excited state of Er 3+ . The authors concluded that such materials may be considered as a promising candidate for Erbium-Doped-Fiber-Amplier (EDFA) operating within the telecommunication window. C. Koepke et al. 24 prepared bers with the same composition doped with 0.6YbF 3 and 0.2ErF 3 (mol%). The luminescence efficiency of the ber was two times higher with respect to the untreated ber. Temperature dependence of Er 3+ luminescence at 1530 nm was observed by varying the temperature from 5 to 350 K. Such dependence was different for the GC ber and associated to a decrease of the phonon energy in the Er 3+ surrounding. In 2012, one of the authors D. Dorosz et al. 25 reported the preparation of Nd 3+ -doped OxGC bers in the system SiO 2 -Al 2 O 3 -ZnO-Na 2 O-SrF 2 using the rod-in-tube method. SrF 2 NCs with size $60 nm were observed aer heat treatment at 635 C. Aer excitation of Nd 3+ ions at 808 nm only the 4 F 3/2 -4 I 11/2 transition was observed, instead the three emissions 4 F 3/2 -4 I 9/2 , 4 F 3/2 -4 I 11/2 and 4 F 3/2 -4 I 15/2 were detected in the bulk GC sample. Moreover, a red-shi of the 4 F 3/2 -4 I 11/2 emission was observed passing from bulk (1060 nm) to ber (1066 nm) and the authors explained this result by an amplied spontaneous emission occurring in the ber. Recently, V. K. Krishnaiah et al. 26 described the preparation and properties of Yb-doped single-index GC bers produced by two methods: (1) by conventional drawing from a glass preform and, (2) by single crucible method (direct-melt process) followed by controlled heat-treatment. The glass system chosen by the authors, a variant of the initial Wang and Ohwaki's composition was: . The authors observed that by using the second method, better quality GCs were obtained with NCs $10 nm homogeneously distributed along the ber. Aer crystallization, higher Yb 3+ PL intensity and quantum yield (0.95) were obtained with respect to the untreated ber (0.82). Some of the authors obtained LaF 3 -GC bers doped with Nd 3+ and demonstrated the possibility to reproduce crystal-like optical features in GC bers. 27 An extended structural characterization was carried out to estimate the dopants distribution and the nanostructure of the GC ber. Other authors obtained uoride NCs in oxide matrices by the melt-in-tube method using borosilicates cladding and drawing temperatures at which the cladding glass soens but the core is completely melted. [28][29][30][31][32][33] For example, Kang et al. 28 obtained NaYF 4 NCs in borosilicate glasses aer heat treatment at 470-500 C and much better ET from Ho 3+ to Er 3+ was observed in the glass-ceramic bers. However, to the best of our knowledge, this method has not still been applied to aluminosilicate compositions because they require the use of SiO 2 cladding and extremely high melting temperatures (1800 C) that can cause strong uorine loss and diffusion phenomena. Other preparation methods in addition to the GC route were also studied, for example Shahzad et al. 34 prepared PMMA bers containing LiYF 4 NCs co-doped with Yb 3+ -Er 3 for applications as optical sensors.
In this study, we report the preparation and upconversion (UC) emission of novel OxGC bers based on cubic solid solutions of the type Na x Lu 2xÀ1 F 7xÀ3 doped with 0.5ErF 3 and 2YbF 3 (mol%). 15 This crystal phase was selected because it shows efficient UC process. [35][36][37][38] Moreover, the Er 3+ /Yb 3+ ratio 0.5/2 showed much better UC emission than Er 3+ -single doped glass. 38,39 The dependence of the upconversion emission on the excitation power intensity is investigated.

Materials preparation
Glasses of composition (mol%) 70SiO 2 -5Al 2 O 3 -2AlF 3 -2Na 2 O-18NaF 3 -3Lu 2 O 3 doped with 0.5ErF 3 and 2YbF 3 were prepared as described elsewhere. 38 Glass rods were polished and used as core materials for the rod-in-tube method. The drawing process was performed at 1230 C and the subsequent crystallization at 600 C for 20 h using a heating rate of 10 C min À1 and a cooling rate of 1 C min À1 to minimize the mechanical stresses between cladding and core. The cross-section of the bers was observed using a ZEISS Axiophot microscope equipped with a Zeiss Axi-oCam MR 5 camera and the ZEN 2.3 lite soware for calibrating the images.

Structural and optical characterization
Powder X-Ray-Diffraction (XRD) was performed using the Bruker D8 Advanced diffractometer (l ¼ 1.54056Å -CuKa1). The powders were milled to a particle size less than 60 mm and both the as drawn bers and GC bers were measured. The diffractograms were acquired in the range 10-60 using a step size of 0.02 . The crystal phase was analyzed with the soware EVADiffractPLUS. The crystallite size, f, was calculated using the Scherrer's equation: 40 where q is the angle of the diffraction maximum, B its full width at half maximum, B i the instrumental broadening and l is the wavelength. The q and B parameters were obtained by tting the peaks to pseudo-Voigt functions. Fourier Transform Infrared Spectroscopy (FTIR) was performed in the range 600-200 cm À1 using a Nicolet 6700-Thermo Scientic in ATR conguration. The resolution was set to 4 cm À1 and 1024 spectra were acquired for each sample.
High-resolution transmission electron microscopy (HRTEM) was performed on glass and GC bers using a JEOL 2100 eld gun emission TEM with a point resolution of 0.19 nm. Selected area diffraction patterns (SAED) were also acquired. Powder samples were dispersed in ethanol and then few drops deposited onto carbon-coated copper grids and the solvent was removed by drying under UV lamp.
Steady-state emission spectra were recorded by transversely exciting the bers with a continuous wave (CW) 980 nm semiconductor laser diode. The uorescence was analyzed with a 0.25 m monochromator, and the signal was detected with a Hamamatsu R636 photomultiplier and nally amplied using a standard lock-in technique.
The transmission losses at 633 nm were determined by using a conventional cut-back (or differential) method. 41 The input for the power was provided by a He-Ne laser and the transmitted power was measured using a Si detector. Then, the ber is cut back and the transmitted power measured again. The loss of the ber in decibels per unit length is obtained from the measured transmitted power ratio of the two measurements.

Glass and GC bers
A detailed thermal characterization of the precursor glass was reported previously. 38 The glass transition temperature T g is around 585 C, the coefficient of thermal expansion (CTE) a around 7.5 Â 10 À6 C À1 , the refractive index is 1.49 at 588 nm and the log h ¼ 4.1 at 1000 C, being h the viscosity in dPa s. DURAN® glass was used as cladding glass, despite T g is around 525 C and CTE is 3.3 Â 10 À6 C À1 , due to its quite high working point (log h ¼ 4 at 1260 C, h in dPa s) and low refractive index (1.473 at 587.6 nm). A stable process was obtained by drawing the bers in the range 1100-1230 C because these temperatures are suitable for the DURAN® cladding glass. However, transparent bers were only obtained by raising the temperature to 1230 C. In fact, oxyuoride compositions tend to crystallize when heated above the T g and as known,the crystallization tendency decreases with faster cooling rates. In the standard procedure as the rod-in-tube method, the cooling of the core is slow down by the thermal screening of the cladding glass. This explains why bers produced with direct-melt are transparent even for drawing at $1000 C while drawing at the same temperature using a cladding glass can cause the spontaneous crystallization of the core. The drawing speed was adjusted to obtain a cladding/core diameter of approx. 225/100 mm as shown in Fig. 1. Fig. 2 shows XRD results of as drawn bers and GC bers treated at 600 C for 10 and 20 h. In the rst case, an amorphous pattern is obtained conrming the absence of crystals in the drawn bers. However, the presence of tiny crystals smaller than 2-3 nm cannot be excluded. For the GC sample, well-dened diffraction maxima are identied and compared to those obtained for the bulk sample treated at the same temperature. It can be concluded that the same cubic solid-solution NaLuF 4 with uorite-type structure is obtained. The reference pattern is the one of NaLuF 4 (JCPDS 27-0725) which has the same cubic structure. The crystal size aer heat treatment at 600 C for 20 h is $20 nm, while for bulk GC, treated in the same condition, the crystal size is $42 nm. 38 This phenomenon was well studied in a previous paper. 27 Due to higher cooling rate of the drawing process, the initial ber present less pronounced phase separation with respect to a bulk material of the same composition. Since the phase separation zones are the precursor for crystallization, there is a delay in the crystallization process of the bers. Therefore, for the same heat treatment temperature, smaller crystals are formed in GC bers with respect to bulk samples. Another important point to outline is the increase of the glass contribution (halo around 20 ) in the XRD pattern due to the presence of the cladding glass.

FTIR
A further conrmation of NCs growth is obtained by FTIR spectra in the range 600-200 cm À1 given in Fig. 3. For the Fig. 1 Transversal section of the fiber.  This journal is © The Royal Society of Chemistry 2019 untreated ber, the absorption band can be simply deconvoluted in two bands corresponding to Si-O-Si rocking vibrations and skeletal deformation of Si-O-Si links centered at $460 and 411 cm À1 , respectively. 42,43 For the GC ber treated at 600 C for 20 h, the same Si-O-Si vibrations are obtained but a further contribution, that can be represented by a Gaussian centered at $330 cm À1 , broadens the spectrum and is associated to the growth of uoride NaLuF 4 -type NCs. 44 The relevant broadening of this band can be related to the nature of the NCs that is a cubic solid solution which most general formula is Na x Lu 2xÀ1 F 7xÀ3 . 15,38 HRTEM HRTEM micrographs of the as made bers are shown in Fig. 4(a  and b). Typical phase separation droplets are observed in agreement with the behavior of the bulk glass. 38 Aer heat treatment at 600 C for 20 h the presence of uoride NCs is clearly observed, Fig. 4(c and d). A broad crystal size distribution is observed and the data t well with two Gaussian functions centered at 15 and 23 nm. By a weighted average of the two contributions, an average crystal size $20 nm is obtained, conrming the XRD results. Fig. 4(e) shows the SAED pattern for the GC ber treated at 600 C for 20 h, conrming the presence of NCs. The analysis of the SAED pattern and its comparison with the JCPDS pattern of NaLuF 4 has also allowed the assignment of the Miller indexes. The corresponding values for the (111), (220) and (311) planes are 0.30, 0.19 and 0.16 nm, respectively, in good agreement with the value reported for the NaYF 4 but with smaller values due to the smaller size of Lu 3+ , Er 3+ and Yb 3+ with respect to Y 3+ . 45 Optical properties Losses. The propagation losses of these multimode optical GC bers codoped with 0.5% ErF 3 -2 mol% YbF 3 are about 10 dB m À1 . This value is similar to that obtained for different GC optical bers. 28,29,46 UC emission. Room temperature UC emission spectra were performed in GC optical bers codoped with 0.5 mol% ErF 3 and 2 mol%YbF 3 , by exciting at 980 nm in resonance with the 4 I 11/2 (Er 3+ ) and 2 F 5/2 (Yb 3+ ) levels. According to previous results obtained in bulk GC samples, this YbF 3 concentration maximizes the overall Er 3+ UC luminescence. Further addition of Yb 3+ ions leads to no further increase in the Er 3+ UC luminescence probably due to energy migration among Yb 3+ ions and the presence of energy back transfer. 37 The UC emission spectra of the GC bers obtained under 980 nm excitation show the characteristic green emissions, attributed to the transitions from the two thermally 2 H 11/2 and 4 S 3/2 levels to the ground state, together with the red emission corresponding to the 4 F 9/2 / 4 I 15/2 transition. In addition, a weak blue emission is observed at around 410 nm which corresponds to the 2 H 9/2 / 4 I 15/2 transition. As an example Fig. 5 shows the UC emission spectrum for an excitation power density of 19 W cm À2 . The UC emission spectrum of the glass ber (GF), obtained under the same experimental conditions, has also been included for comparison. As can be observed, the spectrum of the GC ber shows better resolved bands together with an enhanced intensity by more than one order of magnitude which suggests the incorporation of the rare-earth ions in the crystalline phase. A uorine-rich environment around the rare-earth ions in the GC ber, with lower maximum phonon energy than in the glass matrix, leads to a reduction in the vibration energy of the phonons coupled to Er 3+ ions which benets the UC emission. Thus, the incorporation of the RE ions in the uoride NCs leads to a reduction of the multiphonon relaxation rates and to an increase of the quantum efficiency of emission of rare-earth ions. Moreover, the ratio between the red and green emission intensities increases from 0.95 in the glass ber to 1.36 in the GC ber. This enhancement of the red emission in the GC ber, previously observed in other glassceramics, may be due to an increase of the energy transfer processes populating the 4 F 9/2 level. This behaviour could be associated to a higher concentration of rare-earth ions in the NCs, which reduces the ion-ion distances and increases the probability of energy transfer processes. 47,48 The presence of the weak UC blue emission due to the 2 H 9/2 / 4 I 15/2 transition in the GC ber is remarkable because it reveals the high efficiency of the GC ber as this emission is generally not observed due to the low efficiency of the three or four photon UC processes.
The red emission is strongly enhanced under 980 nm excitation, which means that other excitation processes, as well as multiphonon relaxation from the 4 S 3/2 level, populate the 4 F 9/2 level. This conclusion is conrmed by the visible emission spectrum obtained under 488 nm excitation, in which the main emission corresponds to the green one. The red emitting level 4 F 9/2 can be populated by an excited state absorption (ESA) from the 4 I 13/2 level populated from the 4 I 11/2 level or by an energy transfer process described by ( 2 F 5/2 / 2 F 7/2 ) (Yb 3+ ):( 4 I 13/2 / 4 F 9/2 ) (ETU2) (Er 3+ ). As it was observed in the bulk samples, the presence of Yb 3+ ions has a signicant inuence, not only on the overall UC emission intensity, but also on the luminescence color through increasing the red emission intensity relative to the green one. This behavior, previously observed in other co-doped systems, has been attributed to an increase of the population of the 4 F 9/2 and 4 I 13/2 levels due to energy back transfer from Er 3+ to Yb 3+ through the ( 4 S 3/2 / 4 I 13/2 ) (Er 3+ ):( 2 F 7/2 / 2 F 5/2 ) (Yb 3+ ) crossrelaxation process. [49][50][51] This process reduces the population of level 4 S 3/2 and consequently the green emission. At the same time, the energy transfer described by ( 2 F 5/2 / 2 F 7/2 ) (Yb 3+ ):( 4 I 13/2 / 4 F 9/2 ) (Er 3+ ) (ETU2) populates the 4 F 9/2 level giving rise to an enhancement of the red emission. The presence of energy back transfer has been conrmed in the bulk samples by the lifetime decrease of the green emission from level 4 S 3/2 (Er 3+ ) obtained under 486 nm excitation. 39 Finally, the weak blue emission from 2 H 9/2 level should be due to a three-photon UC process. There are different UC processes to explain the population of this level. Firstly, an ESA process from the 4 F 9/2 level can occur, promoting the Er 3+ ions to the 2 H 9/2 level or an energy transfer process described by ( 2 F 5/2 / 2 F 7/2 ) (Yb 3+ ):( 4 F 9/2 / 2 H 9/2 ) (Er 3+ ) (ETU3). In this process, once the 4 F 9/2 is populated via ETU2, an Yb 3+ ion also in the excited state transfers its energy to the Er 3+ ion which  This journal is © The Royal Society of Chemistry 2019 is promoted to the 2 H 9/2 state being the excess of energy dissipated by the lattice. 52 In addition, we can not exclude the possibility that the blue UC emission occurs via Er 3+ ions since the blue emission has also been observed in single doped GC bulk samples. 15 In this case, two Er 3+ ions in the 4 S 3/2 level can interact through a nearly-resonant energy transfer process determined by the pair of transitions ( 4 S 3/2 / 4 I 9/2 ):( 4 S 3/2 / 2 H 9/2 ). Another possibility is the quasiresonant ETU mechanism ( 4 S 3/2 / 4 I 11/2 ):( 4 S 3/2 / 2 G 11/2 ). 53 Besides, the 2 H 9/2 can be populated through a process in which the Er 3+ ions in the 4 F 9/2 level relax to the 4 I 13/2 level and can transfer part of their energy to those in the 4 S 3/2 level which are promoted to the 4 G 9/2 according to the pair of transitions ( 4 F 9/2 / 4 I 13/2 ):( 4 S 3/2 / 4 G 9/2 ). 54 Then, the 2 H 9/2 level is reached by multiphonon-relaxation from the upper levels.
Pump power dependence of UC luminescence. To further investigate the excitation mechanisms involved in the UC luminescence in the GC bers aer 980 nm excitation, the UC emission spectra has been obtained at different pump power densities.
In all cases, the spectra are characterized by the blue, green, and red emissions corresponding to the 2 H 9/2 / 4 I 15/2 , ( 2 H 11/2 , 4 S 3/2 ) / 4 I 15/2 , and 4 F 9/2 / 4 I 15/2 transitions respectively, but the relative intensity of the emissions is strongly dependent on the excitation power density. As an example Fig. 7 shows the UC emission spectra of the codoped GC ber obtained at two different pump power densities, 2.5 W cm À2 (unfocused) and 0.75 kW cm À2 (focused). It is noticed that the UC emission exhibits signicant changes in the red-to-green emission ratio and in the blue band intensity. As can be seen in the high power limit z 0.75 kW cm À2 , the blue emission increases signicantly and the red-to-green ratio decreases. Fig. 8 shows the evolution of the red-to-green ratio for two different pumping regimes (a) from 1.3 to 19 W cm À2 (unfocused) and (b) 0.025 to 0.75 kW cm À2 (focused). The red-to-green ratio decreases from z3 to 0.65 when the pump power density changes from 1.3 W cm À2 to 0.75 kW cm À2 which indicates that the emission color changes from yellow to green. The color change is easily detected with the naked eye. A similar behavior is observed in bulk GC samples codoped with 0.5 mol% ErF 3 -2 mol% YbF 3 in which the red-to-green ratio decreases from z3.5 to 0.79 when the pump power density changes from 1.3 W cm À2 to 0.75 kW cm À2 .
The calculated chromaticity coordinates calculated for two laser power densities, 2.5 W cm À2 and 0.75 kW cm À2 , are illustrated in Fig. 9. The stars represent the CIE coordinates. The chromaticity coordinates move from yellow (x ¼ 0.4458, y ¼ 0.5031) to green (x ¼ 0.3361, y ¼ 0.6097) when the laser power density increases from 2.5 W cm À2 to 0.75 kW cm À2 which means that the output color can be tuned by controlling the excitation power density.
The upconverted emissions did not show any remarkable thermal load effect induced by the pump laser at 980 nm within  the power range utilized either in bulk or in the GC ber as evidenced by a nearly constant ( 2 H 11/2 / 4 I 15/2 )/( 4 S 3/2 / 4 I 15/2 ) uorescence intensity ratio. This behavior agrees with those found in other rare-earth-doped glass-ceramics pumped with power densities of the same order. 17 To obtain more information about the UC processes, the dependence of the green and red UC emissions has been obtained as a function of the incident pump power density. It is well known that the UC emission intensity (I up ) depends on the incident pump power (P pump ) according to the relation I up f (P pump ) n , where n is the number of photons involved in the pumping mechanism. When UC emission is excited by sequential absorption and energy transfer UC of n photons, its dependence on the incident pump power P pump decreases from (P pump ) n to P pump as long as the UC rate exceeds the decay rate from the intermediate states. 55,56 Fig . 10 shows the logarithmic plot of the UC green and red emission intensities of the co-doped GC ber as a function of the pump power laser density for two different pumping regimes (1.3 to 19 W cm À2 ) and (0.025 to 0.75 kW cm À2 ). As can be seen for the red emission, the slope decreases from 1.26 to 0.86 as the pump power density increases, which indicates that only one photon is required to excite the electrons to the 4 F 9/2 state. However, it is clear from the energy level diagram of Er 3+ ions that two photons are needed to populate this level. A similar behavior is observed for the green emission for which the slope decreases from 1.58 to 1.03. In the case of the blue emission attributed to the 2 H 9/2 / 4 I 15/2 transition, according to the energy level diagram, this emission should be due to a three-photon UC process. However, as can be seen in Fig. 11, the pump power dependence of this emission measured in the high pumping regime between 0.025 and 0.75 kW cm À2 (the blue emission can be accurately measured only in the high pumping regime), gives a slope of 1.40.
This behavior, previously observed in other systems has been attributed to the competition between the decay rate of the intermediate state and the UC rates. This effect was theoretically described by Pollnau et al. 55 The model predicts that the UC luminescence from a state that requires n excitation photons will have a slope of n in the low-power regime when the luminescence intensity is plotted in a double logarithmic representation versus absorbed pump intensity. Higher pump powers and consequently increasing competition between linear decay and UC for the depletion of the intermediate excited state results in a reduced slope. It was experimentally   observed that when UC dominates over linear decay for the depletion of the intermediate excited state, the slope of the luminescence from the upper state n is almost linear. 55 As suggested by Suyver et al., 56 this saturation effect proves that in these GC bers a sensitized energy transfer upconversion process is responsible for populating the emitting levels of Er 3+ .

Conclusions
Transparent glass-ceramics were successfully obtained by the rodin-tube method and subsequent crystallization process. Cubic NaLuF 4 nanocrystals were obtained aer the heat treatment at 600 C, being the average crystal size around 20 nm. Intense green and red upconversion emissions due to ( 2 H 11/2 , 4 S 3/2 ) / 4 I 15/2 and 4 F 9/2 / 4 I 15/2 transitions respectively together with a more weak blue emission due to 2 H 9/2 / 4 I 15/2 transition of Er 3+ ions has been observed under excitation at 980 nm in sodium lutetium uoride glass-ceramic optical bers codoped with Yb 3+ ions. The intensity of the green and red upconversion bands shows a nearly linear dependence on the excitation power which can be explained by saturation effects in the intermediate energy states. The upconversion emission color changes from yellow to green by increasing the excitation power density which allows the color output of the Er 3+ emission to be tuned. Increasing the excitation power density from 1.3 W cm À2 to 0.75 kW cm À2 , the red-to-green ratio decreases from z3 to 0.65, which indicates that the emission color changes from yellow to green as can be easily recognized by naked eyes. The observed behavior makes these optical glass-ceramic bers promising candidates for photonic applications.

Conflicts of interest
There are no conicts to declare.