Patricia Donnadieu

Souad Benrhaiem

Gilles Renou

Chunyang Zhang

Catherine Tassin

Jean-Jacques Blandin

Deformation of Mg-γMg17Al12 in situ composites: Room temperature mechanical behaviour, microstructures and mechanisms

Keywords: intermetallics, complex metallic alloys, Mg-Al alloys, in situ composite, eutectic, room temperature deformation

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Introduction

In the last decade, new composites containing complex metallic alloy (CMA) particles as reinforcement have been developed to improve the strength of light weight alloys. CMAs are intermetallic compounds having in common specific structural features, namely a local atomic tetrahedral packing leading to large and complex atomic clusters [START_REF]Complex metallic alloys: fundamentals and applications[END_REF]. This includes quasicrystals, metallic glasses as well as the many intermetallic alloys earlier called Frank-Kasper phases [START_REF] Frank | Complex alloy structures regarded as sphere packings. I. Definitions and basic principles[END_REF] or Tetrahedrally Close Packed (TCP) phases [START_REF] Sinha | Topologically close-packed structures of transition metal alloys[END_REF]. While CMA phases are hard and usually brittle as bulk materials, composites containing CMA particles can exhibit high strength and large ductility. For instance, strength of ~ 300 MPa with ductility of ~15 % has been reported for Aluminium-based CMA composites containing large volume fraction of metallic glasses [START_REF] Scudino | Production and mechanical properties of metallic glass-reinforced Al-based metal matrix composites[END_REF] or Al3Mg2 particles [START_REF] Scudino | Powder metallurgy of Al-based metal matrix composites reinforced with β-Al3Mg2 intermetallic particles: Analysis and modeling of mechanical properties[END_REF] or quasicrystalline particles [START_REF] Singh | Compressive strength and yield asymmetry in extruded Mg-Zn-Ho alloys containing quasicrystal phase[END_REF] [START_REF] Kabir | Hot isostatic pressing synthesis and mechanical properties of Al/Al-Cu-Fe composite materials[END_REF]. These composites with a large amount of CMA particles (i.e. volume fraction up to 0.4 and even more) were prepared by powder metallurgy. But CMA composites can be also prepared by solid state phase transformation like precipitation [START_REF] Pramod | Aluminum-Based Cast In Situ Composites: A Review[END_REF], [START_REF] Aikin | The Mechanical Properties of In-Situ Composites[END_REF]. However as illustrated by Al-based composites containing quasicrystals, this process usually leads to a limited volume fraction of precipitates [START_REF] Singh | Compressive strength and yield asymmetry in extruded Mg-Zn-Ho alloys containing quasicrystal phase[END_REF]. On the other hand, when the phase diagram of the system displays a eutectic mixture made of a simple and a complex metallic alloy, solidification in the vicinity of the eutectic point can be very efficient to prepare in situ composites with various volume fractions of complex phases. Among many possible eutectic systems, the binary Mg-Al system [START_REF]Pandat-Phase Diagram Calculation Engine for Multicomponent Systems[END_REF] is of particular interest since its phase diagram displays a eutectic equilibrium involving the magnesium-rich solid solution (Mg) and the Mg17Al12 phase that is one of the simplest CMA phases: namely a crystalline cubic structure with cell parameter a=1.04 nm and "only" 58 atoms per cell [START_REF] Villars | Pearson's Crystal Data -Crystal Structure Database for Inorganic Compounds[END_REF]. In a previous work , a series of Mg-Mg17Al12 composites containing a variable volume fraction of (Mg) dendrites and a eutectic mixture composed of (Mg) and phase were prepared by solidification [START_REF] Donnadieu | Preparation, microstructure and properties of magnesium-γMg17Al12 complex metallic alloy in situ composites[END_REF]. These Mg- composites are characterized by a remarkable feature: the magnesium soft phase is embedded in the hard and brittle phase. This microstructure strongly reminds that of the recently developed Bulk Metallic Glasses (BMG) in situ composites [START_REF] Hays | Microstructure controlled shear band pattern formation and enhanced plasticity of bulk metallic glasses containing in situ formed ductile phase dendrite dispersions[END_REF], [START_REF] Qiao | In-situ Dendrite/Metallic Glass Matrix Composites: A Review[END_REF]. The remarkable high strength and high ductility exhibited by the BMG composites are attributed to the presence of soft dendrites embedded in the BMG matrix. Therefore it seems relevant to study the mechanical behaviour of the Mg- composites paying a particular interest to the microstructure evolution with deformation.

The present work will report on the mechanical behaviour of the series of Mg- composites through compressive tests at room temperature, the related microstructures investigated by Scanning Electron Microscopy (SEM) and Transmission Electron Microscopy (TEM) and the interpretation of the results in terms of deformation mechanisms of the Mg- composites with emphasis on the  phase structural specificities.

Materials and methods

Magnesium-based composites with various Mg17Al12 phase volume fractions (60%, 50%, 40%, 30%, 20%) have been prepared by levitation melting from commercially pure ingots of magnesium and aluminium. For information, the (Mg)- eutectic mixture corresponds to 64.5 vol.% of the phase. The experimental protocol to obtain composites with the desired volume fraction is described in details in [START_REF] Donnadieu | Preparation, microstructure and properties of magnesium-γMg17Al12 complex metallic alloy in situ composites[END_REF] . The Mg- composites will be labelled Mg-x% where x% refers to the phase volume percentage.

Table 1 displays for each targeted composite (line 1), the composition measured by inductively coupled plasma atomic emission spectroscopy (ICP-AES) (line 2), the volume fraction of the (Mg) dendrites measured by image analysis (line 3) and the corresponding eutectic mixture volume fraction (line 4). In line 5, the phase volume fraction is calculated from the eutectic volume fraction assuming that the phase ratio in the eutectic mixture is 0.645 as expected from the phase diagram computed with the Pandat software and the PanMg8 database [START_REF]Pandat-Phase Diagram Calculation Engine for Multicomponent Systems[END_REF]. For further use, Table 1 also gives the total (Mg) volume fraction, i.e. fibre + dendrite (line 6) and the volume fraction of the (Mg) fibres in the eutectic mixture (line 7). Compressive tests were carried out at room temperature using a MTS-4M machine operating at a strain rate of 2.5x10 -4 s -1 . Samples for mechanical testing were machined by electro erosion as cylinders (diameter 3 mm, length 5 mm). As-cast and deformed samples were characterized by Scanning Electron Microscopy (SEM) on a LEO Stereoscan 440 and on a JEOL JSM-IT500HR LV. SEM images were acquired using the Back-Scattered Electron (BSE) detector. Samples for SEM were mechanically polished with diamond paste down to 0.25 micron or down to 3 microns with 1 micron water free OPS for final polishing. Specimens for TEM were mechanically polished with SiC paper down to a 50 micron thickness (final step at grade P2400) and ion milled with Argon ions accelerated at 4keV under 10° incidence and at 3keV under 3 ° in a GATAN PIPS 691 or at 3 keV under 7° and at 1 keV under 3° in a GATAN PIPS 2.

A JEOL JEM 3010 Transmission Electron Microscope (TEM) operating at 300 kV was used for Bright Field and Dark field imaging. A JEOL JEM 2100F operating at 200 kV equipped with the nanodiffraction tool ASTAR ™ [15] was used to retrieve orientation and phase maps by automated crystallographic data analysis. The method is based on the scanning of a narrow parallel beam on the TEM thin foil with acquisition of the diffraction pattern for each position. As explained in details in [START_REF] Rauch | Automated crystal orientation and phase mapping in TEM[END_REF] and briefly recalled below, the crystallographic information is retrieved by indexation of the diffraction pattern data set by comparison to simulated diffraction patterns using a template matching technique. The simulation being done under kinematical condition, a precession by a small angle (~1°) applied to the electron beam allows to reduce the dynamical effect and consequently to improve the indexation accuracy. The angular resolution is determined by the angular step size used to build the template data base as well as by the quality of the diffraction pattern. The resolution is currently better than 0.5°. The quality of the indexation is measured by a matching index and a reliability coefficient. The reliability is related to the ratio of the matching indexes for the two best solutions and by definition ranges from 0 to 100 (for a unique solution). Note that reliability values above 15 are considered as fully satisfactory [START_REF] Rauch | Automated crystal orientation and phase mapping in TEM[END_REF]. Once indexation is achieved, misorientation profiles, orientation and phase maps can be produced.

Independently from the indexation, the diffraction data set can be used to form images called virtual Bright Field or virtual Dark Field images depending on the position of the virtual aperture. This virtual BF and DF images are analogous to the classical TEM BF and DF images, except that they are formed by a post treatment of the diffraction data set. Their spatial resolution is thus limited by the one of the diffraction data set (i.e. the step size or the probe size when the latter is larger than the step size). The point by point difference of neighbouring diffraction patterns can also be used to build correlation coefficient maps that are particularly efficient to highlight any sort of singularity giving rise to change in the diffraction patterns of neighbouring points: i.e. microstructural details like dislocations, phase and grain boundaries. These maps are also able to underline 3D features like tilted interfaces.

Results

It is worth mentioning that the microstructural features reported in the following have been observed in all the composites regardless of thevolume fraction. However, for full information, the phase volume fraction corresponding to each image is indicated in the figure caption.

Microstructure of the as cast Mg- composites

The typical microstructure of the as cast Mg- composites is illustrated by the SEM images in Fig. 1. In these BSE images, due to the Z contrast, the magnesium-rich solid solution (Mg) is evidenced by a contrast darker than the Mg17Al12 phase and the Mg- eutectic mixture. The (Mg) dendrites are embedded in the eutectic mixture composed of (Mg) fibres surrounded by the phase matrix (Fig. 1(a)). A halo of  phase continuously surrounds the (Mg) dendrites except in a few locations as marked by dotted arrows in Fig. 1(a,b). However these local discontinuities do not allow for connecting neighbouring dendrites or dendrite arms. In Fig. 1(a,b,c) several eutectic cells can be observed. They are separated by eutectic cell boundaries characterized by coarse fibres as illustrated by Fig. 1(c). Fine scale information on the eutectic microstructure has been obtained by TEM (Fig. 2). Unlike the previous SEM images, in the TEM images, the phase is darker than the (Mg) phase because of its higher diffraction contrast. Fig. 2(a,b) are consistent with the longitudinal and transverse views of a fibrous eutectic made of (Mg) fibres embedded in a  phase matrix. The (Mg) fibres have a diameter between ~0.1 m and ~0.3 m depending on area (Fig. 2(a,b)) and they frequently form a rather hexagonal pattern (Fig. 2(b)). The distance between fibres is corresponding to a phase spacing of ~0.2 m between the (Mg) fibres. However as illustrated by Fig. 2(c) the fibrous microstructure is not always regular. In Fig. 2(c), several eutectic cells are evidenced by the diffraction contrast resulting from the difference of orientation of the phase matrix. The eutectic cells are separated by boundaries pointed by black arrows in Fig. 2(c). As observed by SEM, a phase halo of thickness 0.5 m is surrounding the dendrite. The phase halo is continuous and the interface with the dendrite is locally flat except at the intersection with a eutectic cell boundary where the interface is making a cusp (Fig. 2(c)).

According to the colour coding of the orientation map of the as cast sample in Fig. 2(d), the fibre orientation varies over the area. A point to point misorientation profile taken across the phase matrix (line marked by the double sided arrow in Fig. 2(d) shows that the misorientation is smaller than 0.5 degree. The relationships between the phase matrix and various (Mg) fibres have been retrieved and compared to the one reported by Gulberg and Ryum [START_REF] Guldberg | Microstructure and crystallographic orientation relationship in directionally solidified Mg-Mg 17 Al 12-eutectic[END_REF]. Experimental details are given in the supplementary material (Appendix A). No agreement was found with the Gulberg and Ryum relationship. Over the area shown in Fig. 2(d), the orientation varies and does not match with any simple orientation relationship. This crystallographic study being not useful in further comparison with deformed states, it is not reported here but provided as supplementary material. Generally speaking misorientations across the cell boundaries have no specific value. Owing to the orientation map, these misorientations can been measured. For the cell boundary in Fig. 2(e), the angle between eutectic cells is 5°. It corresponds to the smallest observed for the Mg- eutectic system. It is worth noticing that in the area shown in Fig. 2(e), opposite to the one in Fig. 2(d), the fibre/matrix orientation relationship is constant within each eutectic cell. In our previous study [START_REF] Donnadieu | Preparation, microstructure and properties of magnesium-γMg17Al12 complex metallic alloy in situ composites[END_REF], using conventional diffraction, we found that the orientation relationship was in agreement with the Gulberg and Ryum one. In the present study, the orientation maps obtained with ASTAR provide more complete information. In particular they reveal that the orientation relationships can vary over an area or be constant and when constant, it is not necessarily in agreement with the Gulberg and Ryum orientation relationships (see Appendix A of the supplementary material).

(b)) zone axis normal (resp. parallel) to the fibre axis) (Mg-60% composite). (c) Interface between dendrite and eutectic mixture. The black arrows point eutectic cell boundaries recognizable to a change of the matrix diffraction contrast, coarser fibres and to a cusp at the interface (Mg-50% composite). (d) and (e): ASTAR orientation maps acquired with beam size 1.5 nm, step size 2nm, precession angle 1.2°. (d): eutectic area showing a  matrix with a unique orientation attested by the constant green colour over the area and by the profile over a distance of 400 nm (black arrowed line) showing point-to-point misorientation below 0.5° (i.e. smaller than the angular resolution). The colour coding of the fibres indicates that the fibre orientation with respect to the matrix is changing over the area (Mg-30% composite). (e): eutectic area featuring a eutectic cell boundary. According to the profile across the boundary, this eutectic cell boundaries corresponds to a misorientation of 5°

, the smallest we have observed for eutectic cell boundaries in the Mg- system (Mg-50% composite).

Deformation behaviour at room temperature

The stress-strain curves obtained by compression tests at room temperature are reported in Fig. 3(a). The low strain deformation is consistent with an elastic deformation characterized by a bulk moduli ranging from 45 GPa to 60 GPa with increasing  phase volume fraction. For all the composites, the stress-strain curves exhibit a deformation stage characterized by an inflexion point at about 100 MPa (Fig. 3(a)). As shown by the magnification in Fig. 3(b), for all the Mg- composites, the inflexion point at ~100 MPa corresponds to a quite similar amount of plastic deformation (~ 0.01). After the low stress deformation stage, one observes a significant strain hardening effect, all the more pronounced with increasing amount of the phase. Composites with 20 % and 30 % phase volume fraction exhibit a significant plastic deformation (up to = 0.1 and 0.05 respectively) with a high maximum stress (= 350 and 450 MPa respectively). On the other hand, the composites with a  phase volume fraction larger than 40% do not show a significant deformation -besides the low stress deformation stage -and break for  < 0.02 at a high stress (> 400 MPa). It can be noted that this transition between a brittle and a ductile behaviour is observed at 40%  phase, i.e. a dendrite volume ratio of 50% (Table 1)

Microstructure of deformed Mg- composites

The microstructures after deformation have been investigated in two states: specimens deformed in the low stress deformation stage and specimens deformed to fracture.

The low stress deformation stage (~ 100 MPa,  ~0.01)

The microstructure of specimens after the low stress deformation stage is shown in Fig. 4. The (Mg) dendrites display traces consistent with thin twins in pyramidal planes (Fig. 4(a)) as well as dislocations lying in the basal and pyramidal planes (Fig. 4(b)). Regarding the eutectic mixture, only a few defect between fibres can be noticed like the one pointed by white dotted arrows in Fig. 4(c). However such defects are already present in the as cast material. TEM images in Fig. 6 display microstructures observed in the magnesium dendrites of a specimen deformed to rupture. The dendrites show strong deformation contrast due mostly to dislocations (Fig. 6(a)). However less typical microstructures of possible relevance regarding the propagation of deformation have been noticed. For instance, Fig. 6(b) shows a deformation contrast band crossing a dendrite as well as the phase halo and continuing in the (Mg) fibres of the eutectic cell boundary. In addition, in samples deformed to rupture, traces of dislocations have been observed in the (Mg) fibres and in the phase matrix of the eutectic mixture (Fig. 7). More precisely the DF TEM image in Fig. 7 In the course of the TEM investigation, it was noted that the contrast of the phase defects was often differing from the usual contrast of dislocations: see for instance the striated contrast in the  phase (Fig. 7(a)) compared to the oscillating contrast of dislocations in a (Mg) dendrite (Fig. 4(c)). This effect can be explained by the complementary images presented in Fig. 8. For instance, in the pile-up observed in the spacing between fibres and the eutectic/dendrite interface (Fig. 8 The diffraction contrast of the eutectic mixture in Fig. 9(a) suggesting that the eutectic cell is fragmented in small domains with different orientation, therefore an ASTAR study has been carried out on similar eutectic area reported in Fig. 9(b,c). The virtual dark field image (Fig. 9(b)) reveals several domains while the orientation maps attest by colour change for small misorientation (Fig. 9(c)). The profiles taken across several domain boundaries (Fig. 9(c)) measure misorientations between 2° and 3.5 ° (Fig. 9(d)). As for the defects with the striated contrast (Fig. 8(b)), these misorientations can be interpreted by a pile-up of dislocations. Assuming edge dislocation with 1/2<111> Burgers vector, the 2° to 3.5 ° would correspond to distance between dislocations of about 10 to 7 nm. In other words, a strong localization of the deformation occurs and the formation of slip bands give rise to domains in the eutectic mixture cells.

It is worth noting that the boundaries observed in the deformed samples (Fig. 9(a), 9(b)) cannot be mistaken with the eutectic cell boundaries. Eutectic cells are several microns large while the domains observed in the deformed samples are smaller than 0.5 micron. Also the misorientations between domains in the deformed sample are significantly smaller than the 5 ° value that is the smallest misorientation observed for eutectic cell boundaries (Fig. 2(e)). Regarding the vicinity of dendrites, Fig. 10 (a,b,c) displays the TEM BF image, the orientation map and the correlation map of an area showing the eutectic mixture and part of a dendrite. Like in the as cast state (Fig. 2 (d)), the colour coding of the orientation map (Fig. 10(b)) points that the (Mg) fibre orientation with respect to the phase matrix slightly varies across the eutectic mixture. In this area the orientation relations are found in partial agreement with the Gulberg and Ryum one (details in Appendix A of the supplementary material). The halo at the dendrite /eutectic mixture interface is indexed as a phase with a high reliability. However, opposite to what is observed in the as cast state, the dendrite/eutectic mixture interface is not flat but displays a series of cusps as marked by arrows on the orientation map (Fig. 10 (b)). The correlation map (Fig. 10(c)) reveals that the (Mg) fibres are characterized by a high density of defects while the phase matrix appears quite featureless except for localized singularities in the vicinity of the cusps mentioned above, i.e. in the phase halo narrow spacing of about 0.5 micron between the dendrite interface and the closest fibre. A series of orientation profiles P1 to P5 has been measured (Fig. 10(c)). For reference a profile (P) has been taken in a featureless area: it is characterized by very small misorientations, less than 0.5° over a 200 nm range as was observed in the as cast state (Fig. 2(d)). On the contrary, the profiles P1 to P5 all display an abrupt change of about 1° to 1.5°. The virtual dark field images shown in Fig. 10(d) suggest that dislocations are present in these narrow spacing. If we assume a pile-up of 1/2 <111> phase edge dislocations, the ~1°-1.5 °misorientation corresponds to a dislocation spacing of 50 nm and a piling of about fifteen dislocations in this narrow area. Actually the virtual DF image features, the localized singularities revealed by the correlation maps and the measured misorientation are consistent with deformation localized in the narrow spacing at the interface with the (Mg) dendrite. We note that this localisation is associated to an accumulation of dislocations in the (Mg) dendrite in the vicinity of a cusp that is marked by a white dotted arrow on the virtual DF field image (Fig. 10(e)).

Discussion

The discussion of the mechanical behaviour of the Mg- composites will examine first what the present study is telling about the phase deformation, then we will analyse the results in terms of deformation mechanisms. Finally, the key effects identified through the Mg- composite study will be discussed with respect to the design of high strength composites.

The experimental section has pointed out the following points concerning the phase plasticity. First the macroscopic behavior and the microstructural investigation of the deformed Mg- composites have evidenced the transmission of deformation through the phase. This transmission is demonstrated by the existence of a macroscopic strain even though the ductile phase ((Mg) dendrites and fibres) is embedded in the phase matrix.

It is confirmed by the microscopic observations of dislocations and twins attesting for the plasticity of the (Mg) dendrites while they are enclosed in the phase. More precisely the microstructural investigation points out the presence of a halo of phase separating the (Mg) dendrites from the closest (Mg) fibres of the eutectic mixture. This phase halo systematically observed in the Mg- composites is due to specific thermodynamic and kinetic features of solidification [START_REF] Donnadieu | Preparation, microstructure and properties of magnesium-γMg17Al12 complex metallic alloy in situ composites[END_REF]. This characteristic feature is then preventing from any connection between the dendrites and makes necessary the transmission of deformation across the Mg- eutectic mixture. As a consequence, the  phase forming the eutectic mixture matrix must show some plasticity. The bulk  phase is known to be hard and brittle at room temperature and to become plastic only above 300°C ( [START_REF] Ragani | High-temperature deformation of the γ-Mg17Al12 complex metallic alloy[END_REF], [START_REF] Mathur | Deformation in the γ-Mg 17 Al 12 phase at 25-278[END_REF], [START_REF] Hagihara | Plastic deformation behavior and operative slip systems in Mg17Al12 single crystals[END_REF]), therefore the plasticity attested by the presence of dislocations reveals a size reduction effect resulting from the submicronic scale of the eutectic microstructure. Indeed such size effect on the plasticity of the  phase was already suggested by nanomechanical studies [START_REF] Zhang | Nanomechanics of Mg-Al intermetallic compounds[END_REF] The fact that the deformation is transmitted between the  phase and the (Mg) fibres raises the more general question of the transmission of deformation between a simple structure (here magnesium) and a structurally complex intermetallic (here the phase). Our investigations have shown that dislocations propagate through the  matrix and the (Mg) fibres without displaying a significant deviation of the slip traces neither a high TEM contrast at the Mg/ interface. Slip transfer in multiphase alloys is considered as obeying to the following criterion: minimization of the angle between the intersection of the grain boundary with the glide planes active in each phase, minimization of the Burgers vector of the residual dislocation left at the grain boundary [START_REF] Misra | Slip transfer and dislocation nucleation processes in multiphase ordered Ni-Fe-Al alloys[END_REF]. Therefore dislocations crossing the Mg/ interfaces without deviation and absence of residual dislocation at the Mg/ interface suggest a high structural proximity between the Mg and  phases. This cannot result from a specific orientation relationship between the (Mg) fibre and the  matrix since the orientation maps have shown that the orientation is not well defined and even slightly varies from fibre to fibre.

From a general point of view, Frank and Kasper have demonstrated that the local tetrahedral atom packing induces structural properties that complex metallic alloys have in common. For instance, while the complex metallic alloys do not show dense atomic planes like the hcp or fcc structure, for a small corrugation of about ± 0.1 nm, they feature dense corrugated layers equivalent of compact dense planes. This is illustrated in Fig. 11 on the phasestructure by a series of successive atomic layers and the corrugated dense layer that can be built in the direction <001> (other directions are presented in the Appendix B of supplementary material). Frank and Kasper have also shown that the coordination polyhedron in complex metallic alloys could be only an icosahedron or three other polyhedra related to the icosahedron. These triangulated polyhedra are associated to the tetrahedral atomic packing. The  phase displays such polyhedral while the Mg hcp structure does not (see supplementary material Appendix C). However it can be verified for hcp Mg and Mg17Al12 phase structure that around each atom, similar atomic distances can be found (see supplementary material Appendix C). This local feature together with the existence of dense corrugated layers in various directions can explain the observations of slip traces traversing the Mg and  phases without deviation.

Though the slip transfer observed between the Mg and  phase can be related to well-known structural specificities of complex metallic alloys, to our knowledge, such property has never been reported yet. This is certainly due to the brittleness of the complex alloy phases that usually prevents from deformation studies at room temperature. Here it seems that a size effect resulting from the composite microstructure has allowed for deformation of the  phase at room temperature providing information that otherwise could be reached only by advanced nanomechanics techniques.

Regarding now the deformation mechanisms, the stress-strain curves suggest to distinguish three stages (stages A, B, C in Fig. 12). The first part of the curve (stage A) corresponds to an elastic deformation, with bulk moduli ranging from 45 GPa for composites with a low volume fraction of  phase to 60 GPa for the highest  volume fractions. These values are respectively close to the elastic moduli of magnesium (about 45 GPa depending on the alloy composition) and to the one of the phase (65 GPa) [START_REF] Zhang | Nanomechanics of Mg-Al intermetallic compounds[END_REF]. The stage B corresponds to a low stress plastic deformation almost forming a plateau for a yield stress of about 100 MPa and showing a quite similar amount of plastic deformation (~ 0.01) for all composites. The fact that magnesium alloys are known for yielding at about 100 MPa is not sufficient to account for this low stress deformation because the (Mg) fibres and dendrites are embedded in the phase matrix that is not expected to yield at low stress. However on deformed sample we have noticed deformation contrast bands traversing a dendrite and continuing along the eutectic cell boundary. This observation suggests that deformation can be transmitted between dendrites along paths following the eutectic cell boundaries. The eutectic cell boundaries being native discontinuities formed during the eutectic solidification, they do not require to be nucleated. Moreover, if the eutectic cell boundaries were yielding at a stress similar to the magnesium one, it could explain why all the composites exhibit approximately the same deformation amount for stage B. For all these reasons, we propose to consider the paths following the eutectic cell boundaries as a first way to transmit the deformation between dendrites.

Concerning the transmission of deformation within the eutectic matrix, our investigations have pointed out that, in samples deformed to rupture, dislocations are preferentially located in the spacing between adjacent fibres or between a fibre and the interface with a dendrite. In these submicrometric spacing, the dislocations are forming pile-ups. Also the formation of small phase domains separated by misorientations of several degrees suggests the localization of a high density of dislocations forming slip bands. Consequently, the dislocation pile-ups and the slip bands can constitute a deformation path across the eutectic mixture. However, compared to the first path following the eutectic cell boundaries, this second path involves defects that need to be nucleated in the homogeneous phase matrix that should be then characterized by a higher yield stress consistent with the deformation stage C.

For summarizing, based on the two paths deduced from the microstructural observations, we propose to interpret the Mg- composites behaviour as summarized in Fig. 12. After the elastic deformation (stage A), the low stress behaviour (stage B) is accountable to the deformation of dendrites and propagation of the deformation across the eutectic mixture along the eutectic cell boundaries. Regarding the higher stress behaviour (stage C), the deformation is transmitted across the phase along a path formed by slip bands and dislocation pile-ups. The plasticity of the phase of the eutectic mixture is due to a reduced yield stress resulting from the submicronic scale of the eutectic mixture. In the stage C, opposite to the stage B, the deformation behaviour depends on the phase volume fraction. For phase volume fraction smaller than 40 %, the composites are plastic while for fractions equal to or higher than 40%, they break at high stress (~ 400 MPa) without plasticity except the low stress one. According to the TEM observations, dislocations are observed in all the deformed composites in stage C. As a consequence, the  phase being the matrix of the eutectic mixture of all the Mg- composites, its yield stress must be lower than the maximum stress exhibited by the weakest Mg- composite. Indeed the behavior of the Mg-20% composite suggests a value of about 180 MPa that is corresponding to a change in the plastic behavior in the stage C (inflexion point marked by the dotted line on stress curve, Fig. 12(a)). The fact that this value is rather close to the yield value of the Mg dendrites is consistent with the short stage B observed for all composites. Also this rather low value is consistent with a similarity of mechanical behavior of Mg and Mg17Al12 at small scale suggested by nanomechanical investigation [START_REF] Zhang | Nanomechanics of Mg-Al intermetallic compounds[END_REF].

Regarding now the influence of the phase volume fraction in the stage C , it can be analysed with respect to the strain accumulation in the phase matrix resulting from the slip bands and dislocations pile-ups. Such strain accumulation at the interface with dendrites will be all the more easy to relieve that the dendrite volume fraction is high. Therefore when the dendrite fraction is low (i.e. the phase volume fraction is high) the damage accumulation in the eutectic mixture may become too high and then leads to rupture. This interpretation is consistent with the change in a brittle /ductile behaviour occurring at 40 % phase volume fraction (50 % dendrite volume fraction).

Furthermore it was noticed that cracks in the eutectic mixture were stopped at the (Mg) dendrites. This behaviour reminds of fatigue crack propagation in two phase materials where cracks propagating in the hard phase are attracted by the softer one [START_REF] Bouchard | Damage and Discrete Crack Propagation Modelling : Some Results and Challenges For 2d And 3d Congurations[END_REF]. This effect requires some contrast between the elastic properties of the two phases [START_REF] Ming-Yuan | Crack deflection at an interface between dissimilar elastic materials[END_REF] which is indeed happening in the Mg- composites. The fact that dislocations pile-ups and slip bands are observed along directions normal to the Mg/ interface is also in agreement with the attraction effect of soft inclusion. Moreover like rubber particles in polymer acting as "craze terminators" [START_REF] Perkins | Polymer toughness and impact resistance[END_REF], soft inclusion can contribute to delay the fracture. This can explain why the Mg- composites with high volume fraction break at quite high stress (400-450 MPa) and suggests that the soft inclusion effect is a key effect for the deformation of Mg- composites.

The deformation mechanism described above for Mg- composites is quite similar to the one earlier proposed by Zhang et al. for BMG-dendrites composites [START_REF] Zhang | Dendrite size dependence of tensile plasticity of in situ Tibased metallic glass matrix composites[END_REF]. However there are some differences. First of all, the high elastic deformation of the BMG matrix allows the BMG-dendrites composites to accommodate the dendrites plastic deformation over a larger extent (frequently up to 2%). On the other hand for Mg-composites, the eutectic matrix is able to yield at low stress because it is heterogeneous and contains native discontinuities while the BMG dendrites composites have an homogeneous featureless amorphous matrix [START_REF] Qiao | In-situ Dendrite/Metallic Glass Matrix Composites: A Review[END_REF]. Regarding the deformation stage of Mg- composites (stage C), the plasticity of the  phase matrix can be explained by the submicronic scale of the eutectic microstructure. The localisation of dislocations and the formation of slip bands can be accounted for by the difficult propagation of dislocations in the complex atomic structure of the  phase. In other words, our interpretation of the stage C behaviour relies on the specificities of the CMA structure and on the small scale of the eutectic microstructure. In addition, the propagation of deformation is characterized by features consistent with a soft inclusion attraction.

In metallic glass ribbons as well as in BMG systems or their composites, similar observations have not been reported. However nanoscale heterogeneities are frequently invoked in the understanding of glass plasticity [START_REF] Wang | Spatial correlation of elastic heterogeneity tunes the deformation behavior of metallic glasses[END_REF], [START_REF] Lu | Structural Signature of Plasticity Unveiled by Nano-Scale Viscoelastic Contact in a Metallic Glass[END_REF]. For instance, atomic simulations have pointed out a medium range order called soft spots that could explain the glass mechanical behaviour [START_REF] Ding | Soft spots and their structural signature in a metallic glass[END_REF], [START_REF] Wang | Nanometer-scale gradient atomic packing structure surrounding soft spots in metallic glasses[END_REF]. Unfortunately because of the amorphous structure direct evidences for these soft spots are difficult to obtain [START_REF] Tian | Measuring Structural Heterogeneities in Metallic Glasses Using Transmission Electron Microscopy[END_REF]. To that respect, eutectic mixture involving a CMA phase could provide attractive model systems since, as pointed out by the corrugated layer described above, the tetrahedral local order of CMA phases makes them quite isotropic and from that point of view close to an amorphous alloy.

The study of the Mg- composites has revealed the combination of three key effects: the structural specificities of CMA phase allowing the transmission of deformation between simple and complex structure, the size effect on the plasticity of the  phase and the soft inclusion effect. This result could be of interest in the design of new high strength intermetallic composites. In the last decades, various systems have been pointed out for their remarkable mechanical properties. Besides the case of BMG dendrites composites, high strength combined to high ductility has been reported for Zr-based composites (the so called MINCs composites) [START_REF] Dusoe | Strong, ductile, and thermally stable Cu-based metal-intermetallic nanostructured composites[END_REF], for light alloys eutectic composites [START_REF] Park | High-strength bulk Al-based bimodal ultrafine eutectic composite with enhanced plasticity[END_REF], [START_REF] Song | Deformation mechanisms of a bimodal eutectic Mg72Cu5Zn23 ultrafine composite[END_REF] and for nanoeutectic alloys [START_REF] Li | FeCo-based multiphase composites with high strength and large plastic deformation[END_REF]. Though most of these systems (if not all) are containing CMA phases, the structural point of view is rarely considered and properties are rather analysed in terms of microstructure scale. However, among the systems listed above, the nanoeutectic composites are very fine scale while the MINCs alloys display coarser features and the light alloys eutectic composites are characterized by a bimodal microstructure. Hence the scale alone does not account for the outstanding properties of these composites. All these systems have also in common to be made of a soft phase and a very hard and often brittle phase which is the matrix. In other words, like in the Mg-Al system, the soft phase is embedded in the hard one.

According to the present study of Mg- composites, such microstructure allows for a soft inclusion effect able to delay the fracture propagation and consequently to reach higher deformation. Though dendrites have been identified as having an important role in the propagation of shear band and cracks [START_REF] Park | High-strength bulk Al-based bimodal ultrafine eutectic composite with enhanced plasticity[END_REF], [START_REF] Park | Improving the Mechanical Properties of Fe-Nb-(Ni-Mn) Dendrite-Ultrafine Eutectic Composites via Controlling the Primary Phase Features[END_REF], the fact that the soft phase was embedded in the hard one and the resulting soft inclusion effect were not specifically emphasized. Indeed the soft inclusion effect makes possible to overcome the usual brittleness of CMA phases and could constitute a useful guideline in the design of composites with a high level of mechanical properties. Finally, one of the main findings of the present work is that instead of trying to reinforce a soft matrix, softening a hard one appears as a route of interest for designing new materials with enhanced properties.

Conclusion

The present work was centered on the room temperature deformation of Mg-Mg17Al12 composites containing various amounts (between 20 and 60 vol.%) of the complex phase. Combining compression tests with finescale microstructural investigations led to the following key features:

 The as-cast microstructure of the composites consists of (Mg) dendrites surrounded by cells of the fibrous (Mg)-eutectic mixture. It can also be described as a matrix of hard  phase containing dendrites and fibres of the soft metallic (Mg) phase. The same microstructural features are observed after deformation. All the composites exhibit plastic deformation under compression at room temperature. It is characterized by two stages. The first one occurs for all composites at low stress (~ 100 MPa), low strain (< 0.01). The second stage is characterized by a deformation behaviour depending on the  phase volume fraction, namely the composites containing 20 and 30 vol.% of the  phase display a significant plastic deformation with strain hardening while the composites with  phase volume fractions equal to or larger than 40% show only negligible additional deformation before breaking at high stress (> 400 MPa).

 The macroscopic behaviour and the microstructural investigation are consistent with the activation of several deformation mechanisms: twinning and dislocation multiplication in the (Mg) dendrites followed by their propagation across the eutectic mixture throughout the matrix and the (Mg) fibres along the eutectic cell boundaries or along dislocation pileups and slip bands resulting from the localization of the phase dislocations.

 The transmission of the deformation between the simple Mg structure and the complex  phase one that has been here experimentally evidenced can be explained by the specificities of the complex  phase structure ; namely the proximity of the atomic distances and the existence of corrugated layers in many directions acting as dense glide planes. These specificities can also explain why slip transfer between the Mg and  phases is possible while the orientation relationship is not well defined.

 The observations are also revealing features of the Mg- eutectic mixture which are of major interest for the mechanical behaviour: -the plasticity of the hard and brittle phase eutectic matrix that is attesting for a size effect resulting from the fine scale of the microstructure, -a localization of deformation in agreement with a soft inclusion effect.

More generally the combination of soft inclusion effect, size reduction and specific atomic features of complex metallic alloy phase provides a frame of understanding for the plasticity of various recently developed composites like the BMG-dendrites and other composites that all have the specificity to involve complex metallic phases and a soft phase embedded in the hard and brittle one. The transmission of deformation between the simple and complex metallic phases is a point that is rarely addressed though it can be of primary interest for developing new metallic alloy composites. Also the soft inclusion effect appears as a key effect that allows to overcome the brittleness of CMA phases and then can give rise to high strength composite systems with a high ductility.

From the characterization point of view, the present microstructural study illustrates how the ASTAR advanced nanodiffraction method constitutes a valuable tool to quantitatively study complex microstructures and significantly complement the classical TEM imaging. This method has been here particularly efficient to investigate dislocation configurations like pile-ups and slip bands in the composites. It appears as a most appropriate technique to investigate and decipher the microstructure of complex/simple alloys composite systems.

Fig. 1 :

 1 Fig. 1: Microstructure of the as cast Mg- composites (SEM-back scattered electron contrast). (a) (Mg) dendrites embedded in the eutectic mixture.(b) Dendrites and eutectic cells imaged at higher magnification. Note the bright halo continuously enclosing the (Mg) dendrites except for some rare discontinuities marked with dotted arrows in figures (a), (b). The morphological details within the various cell sections are consistent with a fibrous eutectic. (c) Eutectic cell boundary recognisable to the change in fibre orientation and to the presence of coarser fibres. (Mg-40% composites)

Fig. 2 :

 2 Fig. 2: Fine scale microstructure of the as-cast state. (a) and (b): TEM images recorded along different zone axes attesting for a fibrous morphology with (Mg) fibres embedded in a phase matrix and forming an almost hexagonal array. (a) (resp.

Fig. 3 :

 3 Fig. 3: Compressive stress-strain curves of the series of Mg- composites at room temperature (a) and magnification on the low stress part (b). Note the inflexion point on the deformation curves present for all composites at about 100 MPa and a deformation 0.01.

Fig. 4 :

 4 Fig. 4: TEM micrographs of specimens deformed up to the 100 MPa inflexion point. (a) the white arrows point several twins in the (Mg) dendrite (Mg-40%composite) (b) Many dislocations lying along various planes can be observed in the (Mg) dendrite (Mg-60% composite). (c)TEM Dark Field image showing dislocations in the (Mg) dendrite and some rare defects (white dotted arrow) in the phase matrix (Mg-40%composite)

Fig. 5 :

 5 Fig. 5: Microstructure of a specimen deformed to fracture -compression direction parallel to the image width (SEM-back scattered electron contrast), (a) Area close to the fractured surface. More or less open cracks are crossing the eutectic mixture. They are close to the normal to the interface between (Mg) dendrite and eutectic mixture. Most of them are at 45 ° from the compression direction. Note that the black traces within the dendrites are due to polishing scratch. (b) Area at about 1 mm from the fractured surface no cracks are observed. As in the as cast state (Fig. 1), discontinuities of the halo surrounding the dendrite are observed. It is particularly obvious that they are associated to eutectic cell boundaries. (Mg-20%composite)

Fig. 6 :

 6 Fig. 6: Microstructure after deformation up to rupture (TEM Bright Field images) (a) Deformation contrast mostly due to dislocations is observed in the (Mg) dendrites. (b) Deformation contrast band traversing a (Mg) dendrite and extending inside a eutectic cell boundary (Mg-30% composite)

  (a) is consistent with dislocations slipping in the plane perpendicular to the (Mg) fibres and successively crossing thephase matrix and the (Mg) fibres. The BF TEM image in Fig.7(b) also shows a trace in agreement with the propagation of dislocation through thephase matrix and a (Mg) fibre. According to the diffraction pattern indexation, the slip plane associated to the dislocation in Fig.7(b) is a (1-12) phase plane. In both Fig.7(a, b), the traces are consistent with dislocations propagating in a plane normal to the (Mg) fibre axis and to the fibre/phase interface.

Fig. 7 :

 7 Fig. 7: Microstructures after deformation up to rupture. (a) Dark Field TEM image of a eutectic area showing fibres with their axis in the observation plane. Slip traces are observed running through the (Mg) fibres and the phase matrix along a direction normal to the Mg/interface. Note the striated contrast exhibited by the dislocation in the phase. (Mg-60% composite) (b) Bright Field TEM image showing traces crossing the phase and continuing in a (Mg) fibre. On the right side, the corresponding  phase matrix diffraction pattern indicates that the trace in Fig. 7(b) corresponds to a (1-12) plane. (Mg-40%composite)

  (a)),  phase dislocations are separated by about 100 nm and they show the usual dislocation oscillating contrast. On the other hand, within the eutectic mixture (Fig.8(b)), traces of defects joining two fibres are showing the anomalous striated contrast earlier mentioned (Fig.7(a)). Actually the corresponding diffraction pattern (Fig.8(b)) shows spot splitting revealing a local misorientation of about 1°. Such misorientations are known for giving rise to a moiré contrast what can account for the striation on the defect traces. Consistently no misorientation was seen on the area showing the dislocation pile-up (Fig.8(a)). Moreover, if we consider a pile-up of  phase edge dislocations with Burgers vectors ½<111> separated by 100 nm, the resulting misorientation is smaller than 0.3°, therefore difficult to detect. Following the same line, the ~ 1° misorientation observed for the area in Fig.8(b) would correspond to a pile-up of dislocations separated by distances smaller than 30 nm. This short separation added to the moiré contrast can explain why dislocations do not show their usual contrast.

Fig. 8 :

 8 Fig. 8: Microstructure after deformation to rupture (TEM dark field images). (a) Dislocations in the phase pile-up in the spacing between the dendrite interface and a close (Mg) fibre. Note that the phase dislocations are isolated and show the usual dislocation oscillating contrast. (Mg-30%composite) (b) Area showing defects between (Mg) fibres characterized by a striated contrast. A local misorientation is revealed by the spot splitting on the diffraction patterns (Mg-50%composite)

Fig. 9 :

 9 Fig. 9: Microstructure after deformation to rupture. (a) TEM bright field images showing micro-sized domains inside the eutectic cell. (b) Virtual Dark Field image revealing domains with different diffraction contrasts separated by boundaries showing dislocation like features. (c) Orientation map of an area located in the eutectic mixture. The colour coding points out small orientation changes in the phase matrix while the fibres exhibit a unique orientation confirming the presence of only one eutectic cell in this area (d) Misorientation profiles across the boundaries labelled P1 to P4. (Mg-50%composite)

Fig. 10 :

 10 Fig. 10: Microstructure after deformation to rupture. (a) TEM bright field image of an area showing the interface between the eutectic mixture and a dendrite. (b) Orientation map obtained with ASTAR The  phase halo is continuous but the interface with the dendrite is not flat. Along the interface, black arrows are pointing a series of cusps. (c) Correlation map showing traces of dislocations in the (Mg) dendrite and a phase matrix free of defects except for singularities between the interface and the closest fibres. The blue arrows point the locations of the misorientation profiles displayed below the correlation map. The reference profile (P) drawn in an area free of defects shows negligible misorientation (< 0.5°, i.e. resolution limit of the indexation). The profiles P1 to P5 are all characterized by an abrupt change of about 1°-1.5°. (d) Virtual Dark Field image revealing features consistent with dislocations in the narrow spacing between the interface and the closest fibres. (e) Virtual Dark Field image showing dislocation traces in the (Mg) dendrite and fibres. The dotted arrow indicates an accumulation of dislocations in the dendrite in the vicinity of a cusp. (Mg-50%composite)

Fig. 11 :

 11 Fig. 11:Mg17Al12 phase structure : atomic planar layers and dense corrugated layer. (a) Perspective view of the Mg17Al12 phase structure cell (full scale radii). (b) Side view of the Mg17Al12 cell corresponding to a projection along the [001] direction (20% radii size). The dotted lines labelled 1 to 6 and the grey rectangle respectively indicate successive atomic planar layers and the selection used to build a dense corrugated layer. (c) Projected and perspective view of the corrugated layer corresponding to the grey rectangle selection. Note that for this selection with 0.23 nm thickness (corrugation ± 0.115 nm), the layer is dense (full scale radii). (d) Atomic planar layers corresponding to the layers 1 to 6. Note that they are all characterized by a low atomic density (full scale radii).

Fig. 12 :

 12 Fig. 12: Deformation behaviour and proposed mechanisms. (a) Stress-strain curves of the Mg- composites revealing three deformation stages labelled A, B, C. Stage A corresponds to the elastic deformation, stages B and C to plastic deformation. The dotted line points out a change in the stage C plastic behaviour of the Mg-20% composites suggesting an estimate for the yield point of the  phase in the eutectic mixture. (b) Stage B : Deformation occurs in the (Mg) dendrites and is transmitted by the eutectic cell boundaries. (c) Stage C : deformation occurs in the eutectic cells with the formation of dislocation pile-ups and slip bands.

Table 1 :

 1 Alloy composition and volume fraction of the various microstructural constituents of the Mg- composites

	Composites	Mg-60% Mg-50% Mg-40% Mg-30% Mg-20%
	Composition (Mg wt. %)	68.1	71.0	74.5	78.9	83.7
	(Mg) dendrites volume fraction	0.07	0.21	0.40	0.56	0.72
	Eutectic mixture volume fraction	0.93	0.79	0.60	0.44	0.28
	 phase volume fraction	0.60	0.51	0.39	0.28	0.18
	(Mg) volume fraction (fibres + dendrites)	0.40	0.49	0.61	0.72	0.82
	(Mg) fibres volume fraction	0.33	0.28	0.21	0.16	0.10

Acknowledgments: This research has benefited from the characterization equipments of the Grenoble INP -CMTC platform supported by the Centre of Excellence of Multifunctional Architectured Materials "CEMAM" n°ANR-10-LABX-44-01 funded by the Investments for the Future programme


A E Souza

S G Antonio

Sidney J L Ribeiro

Douglas Faza Franco

Gustavo Galeani

Thierry Cardinal

Marc Dussauze

Marcelo Nalin

Heavy metal oxide glass-ceramics containing

Keywords: Glass-ceramics, Glass matrix composites, Gallium garnet, Luminescence, Rare earths

come

Introduction

Glasses containing rare earth ions have been studied as emitters in optical amplifiers [START_REF] Yoshimoto | Low phonon energies and wideband optical windows of La 2 O 3 -Ga 2 O 3 glasses prepared using an aerodynamic levitation technique[END_REF] , optical thermometers based on upconversion effects [START_REF] Manzani | A portable luminescent thermometer based on green up-conversion emission of Er 3+ /Yb 3+ co-doped tellurite glass[END_REF] , white light generation [START_REF] Li | Transparent Ce:Y3Al5O12 ceramic phosphors for white light-emitting diodes[END_REF], [START_REF] Hirao | Active glass for photonic devices: photoinduced structures and their application[END_REF] among other applications. In particular, glass-ceramics prepared from glasses by controlled thermal treatment allows to explore properties like those of single crystals. The growth of optically active crystals into a glassy matrix however is a challenge, because the homogeneity and transparency of glass-ceramics have to be sustained while luminescent ions are incorporated in the crystal phase. Several glasses have been used for such purposes, but each one has its own peculiarities. Moreover, cubic phases are preferred for avoiding anisotropy induced by preferential orientations of crystals into the glassy matrix. In this work, heavy metal oxide glasses were studied due to their set of interesting properties, such as low phonon energy, good infrared transparency, high refractive index (higher than 2), good chemical stability, among others, making them excellent as photonics.

Recently, several papers have disclosed the interesting properties of RE 3 Ga 5 O 12 garnets (RE = rare earths) [5], [6] . Almost all rare earths can be used to explore both optical [START_REF] Venkatramu | Synthesis, structure and luminescence of Er 3+ doped Y 3 Ga 5 O 12 nanogarnets[END_REF] and magnetic properties of these garnets P a g e | 2 [START_REF] Paddison | Hidden order in spin-liquid Gd 3 Ga 5 O 12[END_REF] . Most of reports described the synthesis of these materials as polycrystalline nanosized powders or sub micrometer single crystals by solid state reactions, solution combustion and sol-gel methods [START_REF] Venkatramu | Synthesis, structure and luminescence of Er 3+ doped Y 3 Ga 5 O 12 nanogarnets[END_REF], [START_REF] Krsmanović | Structural and luminescence investigation on gadolinium gallium garnet nanocrystalline powders prepared by solution combustion synthesis[END_REF], [START_REF] Zorenko | Luminescent and scintillation properties of Lu 3 Al 5 O 12 :Sc single crystal and single crystalline films[END_REF], [START_REF] Kir'yanov | Powerful visible (530-770 nm) luminescence in Yb,Ho:GGG with IR diode pumping[END_REF] .

The synthesis of new glass compositions and the comprehension about the thermodynamic of crystal growth, mainly concerning the kinetic behavior, have allowed for the preparation of special glassceramics. Luminescent crystals embedded in heavy metal oxide glass are of great interest for the development of efficient amplifiers or powerful solid-state lasers [START_REF] Stewen | Yb:YAG thin disk laser with 1 kW output power[END_REF], [START_REF] Honea | High-power dual-rod Yb:YAG laser[END_REF] . A triple-doped glass (Yb 3+ , Er 3+ and Tm 3+ ) was studied herein because the combination of these rare earths has been found to result in white light generation due to a well-known energy transfer upconversion mechanism [START_REF] Zhang | Effects of PbF 2 doping on structure and spectroscopic properties of Ga 2 O 3 -GeO 2 -Bi 2 O 3 -PbO glasses doped with rare earths[END_REF], [START_REF] Yang | Broadband 1.53 μm emission in Er 3+ -doped Ga-Bi-Pb-Ge heavy metal oxide glasses[END_REF] . In this work, we studied how changes in the chemical composition of heavy metal oxide glasses influence the formation of highly luminescent glass-ceramics. Both glasses and glass-ceramics were characterized by thermal analysis, Differential Scanning Calorimetry (DSC), X-ray diffraction, micro-Raman, micro-luminescence and UV-Vis spectroscopies, optical microscopy (OM), scanning electron microscopy (SEM), and energydispersive X-ray spectroscopy (EDS).

Experimental details

The following analytical grade raw materials were used: GeO 2 (99.999%), Ga 2 O 3 (99.99%) and Bi 2 O 3 (99.975%) from Alfa Aesar, PbO (99.0% -Fluka), Tm 2 O 3 (99.9957%), Er 2 O 3 (99.9695%), and Yb 2 O 3 (99.9987%) from REEtec.

Multicomponent heavy metal oxide glasses with molar composition 40PbO

-20GeO 2 -25Bi 2 O 3 -15Ga 2 O 3 and 97(40PbO-20GeO 2 -25Bi 2 O 3 -15Ga 2 O 3 )-2Yb 2 O 3 -0.5-Er 2 O 3 -0.5-Tm 2 O 3 (
in mol%), here called PGBG and PGBG6, respectively, were studied. The glass samples were obtained by a melt-quenching process, in which the raw materials were weighed, homogenized in a speed mixer equipment, transferred to a Pt/Rh crucible, and melted in a conventional furnace. The melts without rare earths were kept at 1200 °C for 40 min for full homogenization, and then poured into a pre-heated steel-brass mold at 350 °C, where they were annealed for 3 h to eliminate internal stresses due to thermal shock. Afterwards, the furnace was turned off and the samples were allowed to cool down to room temperature.

The glass-ceramics were prepared by the addition of rare earths to the glass compositions. To prepare the glass-ceramics, the melts were cooled into the pre-heated steel-brass mold at 350 °C. To obtain the crystal-free glasses containing rare earths, the melts were cooled between two plaques kept at room temperature (fast quenching rate). On the other hand, when the melts were cooled out of the furnace to room temperature, into the crucible, fully crystallized samples were obtained (slow quenching rate).

The bulk glass and glass-ceramic samples were polished using ethyl alcohol and silicon carbide (SiC -Struers) sieves, with decreasing grain size of 800 (20-24 µm), 1200 (13-16 µm), 2400 (7-10 µm), 4000 (<6 µm), and finally with 0.5 and 0.3 µm alumina suspensions for improving their optical quality.

The samples were characterized by X-ray diffraction (XRD) using a Bruker D8 Advance diffractometer with Cu kα radiation (1.54 Å) in the range of 2(θ) from 2° to 70°, Soller slit 2.5°, divergence slit 0.2° and pass of 0.02°. Thermal analysis (DSC) were done in a NETZSCH 404F3 Pegasus equipment in the range from 300 to 600 °C under N 2 atmosphere at a flow rate of 20 mL min -1 and heating rate of 10 °C min -1 using platinum crucibles. The analyses were performed using around 15 mg of sample. The characteristic temperatures: glass transition temperature (T g ) and onset of crystallization (T x ) were determined. The stability parameter (T x -T g ) was used to compare the stability against devitrification between samples. The errors associated with the determinations of Tg and Tx were±2 °C.

Micro-Raman spectra were obtained using a Horiba Jobin Yvon LAB Ram HR spectrometer using a laser with λ = 632.8 nm, slit of 100 µm, and objective of x50WD in the range from 200 to 1500 cm -1 .

Micro-luminescence measurements were recorded in backscattering mode on a micro-Raman spectrometer HR800 (Horiba/Jobin Yvon). As excitation source, a continuous wave laser operating at 638 nm and a pulsed picosecond laser of 1064 nm were used. Typical spectral resolution used for Raman is 2.5 cm -1 . The objective used was a 100x long working distance with a numerical aperture of 0.5 from Mitutoyo allowing a resolution in the X-Y plane of 0.8 µm.

Emission spectra were obtained in the range from 250 to 850 nm using a Horiba Jobin Yvon spectrophotometer equipped with a R928 Hamamatsu photomultiplier and excited by a diode laser operating at 980 nm. The UV-Vis spectra were obtained using a Cary 5000 -Varian spectrometer in the range from 350 to 2000 nm in the transmission mode.

Scanning electron microscopy (FEG-SEM) analyses were done using a JEOL microscope model 7500F. Chemical analysis were obtained using EDS and elemental mapping on bulk samples covered with gold, using a detector ThermoFisher Scientific Ultra Dry with voltage of 10.0 kV and take off angle of 49.1°. Optical images were collected using an optical microscope model Leica DM2500M with lens 5, 10, 20 e 50x. To obtain the images, the glass samples were polished to expose the crystal faces.

Results and discussion

Melt-quenching is the most common technique used to obtain completely amorphous glasses. In this work, this technique was combined with the chemistry of supersaturated solutions. First, the heavy metal oxide glass composition was chosen to obtain samples with high transparency from the visible up to medium infrared region. As a second step, we studied the upper limit rare earth content that could be incorporated in the glass matrix without inducing its complete crystallization. In practical terms, the synthesized glass containing rare earths should be transparent to the eyes, without visible light scattering. In the last step, we studied the synthesis conditions, in specific the difference between the melting and mold temperatures, thus controlling the quenching rate, as well as the annealing times.

Based on the steps described above, the chemical composition 40PbO-20GeO 2 -25Bi 2 O 3 -15Ga 2 O 3 was chosen as the passive host while the composition 97(40PbO-20GeO 2 -25Bi 2 O 3 -15Ga 2 O 3 )-2Yb 2 O 3 -0.5-Er 2 O 3 -0.5Tm 2 O 3 was the active one. The concentration of rare earths (3 mol% in total) was above the limit to obtain the completely amorphous glassy phase, in other words, this concentration represents a supersaturated condition. Both glass and glass-ceramic samples were characterized, as discussed below. Fig. 1a shows photographs of the PGBG and PGBG6 samples. Visually, no macro-differences were observed, and both samples seemed to be transparent. However, optical microscope observations revealed the presence of cubic crystals ranging from 5 to 100 µm in the glass-ceramic sample, Fig. 1b (bottom), while the sample PGBG (top) did not show any trace of crystalline phases within the same scale. In Fig. 1c, it is shown a detailed image of the cubic crystals. This image will be used latter in the text to discuss the luminescence properties of both crystals and glasses.

Fig. 1. A) Photographs of glass (PGBG) and glass-ceramic (PGBG6) samples. B) Optical microscopy images of glass PGBG (top) and glass-ceramics PGBG6 (bottom) with lens x10 (scale bars correspond to 50 µm). C) Magnified OM image showing the cubic crystals in detail.

Fig. 2 shows the DSC curves of the glass and glass-ceramic samples. It is possible to ensure that the glass phase was present in both samples. A Tg was observed at 379 °C for PGBG and no Tx and Tp (temperature of maximum crystallization) were observed in the temperature range analyzed. The PGBG6 sample displayed Tg at 390 °C due to addition of rare earth ions. No evidence of further crystallization was observed up to 600 °C, thus it is possible to affirm that the thermal stability parameter (Tx-Tg) was higher than 200 °C for both samples.

Fig. 2. DSC curves of PGBG and PGBG6. Tg values were obtained as the interception between two tangent dashed lines.

Due to the fact that the glassy phase is not detected by XRD, resulting only in a halo characteristic of amorphous materials (insert in Fig. 3), the discussion about the X-ray diffraction data was done using the glass-ceramic sample. Fig. 3 shows the X-ray diffraction profile of the glass-ceramic sample along with the simulated pattern. It can be observed a halo centered at 29° due to the amorphous phase, and narrow diffraction peaks at 18°, 36° and 56° of 2θ. From the XRD measurements, it was possible to identify the crystal phase using the Rietveld method (TOPAS-Academic v6.0). Data refining confirmed that the obtained phase is due to the crystalline phase of Yb 3 Ga 5 O 12 (PDF-ICDD: 73-1373) with space group Ia-3d.

The amorphous fraction was corrected with the background using the Chebyshev function. The unit cell was refined, and the cell parameter obtained was a = 12.229(2). The preferred orientation was corrected with the March Dollase model. The results were also compared with those of germanium garnet phases. Several crystalline phases based on the generic structure A 3 B 2 Ge 3 O 12 (with A= twofold cations and B= threefold cations) have been reported in the literature [START_REF] Liu | Synthesis, crystal structure, and enhanced luminescence of garnet-type Ca 3 Ga 2 Ge 3 O 12 :Cr 3+ by codoping Bi 3+[END_REF] . Since these compounds can also form cubic structures with the same space group, Ia-3d, it was necessary do discuss them to justify the formation of gallium instead germanium garnets in the samples prepared herein. First, it is worth pointing out that the chemical composition of the PBBG6 glass would allow the Pb 2 Ga 3 Ge 3 O 12 compound to be formed. However, the crystallographic parameters of this phase were not found in the literature. A similar compound with chemical composition Ca 2 Ga 3 Ge 3 O 12 presents the same specifications of gallium garnets, including a lattice parameter P a g e | 6 a = 12.251 Å. This value is close to that obtained in this work (the effective ionic radii of Ca 2+ and Pb 2+ are 100 pm and 119 pm, respectively, thus it is necessary to consider some changes in the a-value). Table 1 summarizes the data of all discussed compounds. It can be observed that all crystals exhibited the same space group and the a-values fluctuated from 12.204 Å for Yb 3 Ga 5 O 12 to 12.255 Å for Er 3 Ga 5 O 12 , depending on the ionic radius of each rare earth.

Due to similarities in packing structure and lattice parameters, it was quite difficulty to confirm the rare earths distribution within the crystals. However, there was strong evidence that all rare earth ions were included both in the glassy and crystalline phases, as shown below.

Fig. 4 represents the 3D unit cell and its packing for the formation of crystalline structures. The RE 3 Ga 5 O 12 structure may be generically described by GaO 6 and GaO 4 polyhedra linked by an oxygen atom, forming 3D chains along the network, and creating dodecahedral cavities occupied by rare earth atoms [START_REF] Venkatramu | Synthesis, structure and luminescence of Er 3+ doped Y 3 Ga 5 O 12 nanogarnets[END_REF] . The distribution of chemical elements in the PGBG6 sample was studied by EDS and SEM elemental mapping, and the results can be seen in Figs. 5 and6, respectively. These measurements were done after polishing the glass-ceramic until one crystal surface was exposed. The EDS results strongly indicated that the chemical composition of crystals mainly contained Ga rather than germanium. The energies of both atoms were similar, however, as shown in detail in Fig. 5, it was possible to identify the presence of Ga at 0.95 (edge L 1 ) and 1.098 keV (edge L a1 ) as the most intense peaks, while the Ge signal at 1.19 keV (edge La 1 ) was only observed as a shoulder. Also, a signal ascribed to Yb 3+ is noted at 1.19 keV (edge Mz I ). This fact supports the formation of gallium garnet structures. The electronic transitions arising from the three rare earths could be identified between 1.3 and 1.6 keV. The signals at 2.33 and 2.42 keV were ascribed to possible transitions of Pb and Bi, respectively. However, such transitions may overlap with those of gold, which was used to coat the sample. It was assumed that both, Pb and Bi, were absent in the samples, once the elemental mapping, shown in Fig. 6, rules out the presence of such ions into the crystals. These results confirm the formation of the gallium garnet phase in the glass-ceramics.

Fig. 6. SEM image and elemental mapping results of PGBG6. The darker part in the SEM image corresponds to the crystal, while the lighter part relates to the glassy phase. The map of each element is displayed along with its chemical symbol.

The darker region in the center of the SEM image corresponds to the crystal, while the lighter region is ascribed to the glass. Fig. 6 reveals the distribution of chemical elements contained in both phases. The dark region is predominantly composed of Ge, Ga, and the three rare earths (Yb, Er and Tm). The presence of Ge and Ga in the crystal is due to the fact that these elements have close atomic numbers and, consequently, their transitions are similar, making difficult to separate both signals. Thus, combining these data with those from Rietveld refinement and EDS, it is possible to confirm that gallium is present in the crystals. The same analogy is done for the rare earth ions, since these elements are neighbors in the periodic table, thereby exhibiting similar transitions. Accordingly, the elemental maps indicated the presence of Yb, Er and Tm. As will be discussed below, the luminescence data indicated that the three REE were incorporated in the crystals. However, we also observed that these elements were dispersed in the entire sample surface because the crystallization occurred only partially, and as a result, a fraction of REE remained in the glassy phase.

In the OM image shown in Fig. 4, the cubic crystals seem to be partially filled, since holes at their surface are noticeable. These features were also observed by SEM. Such a behavior will be explained in detail below using the micro-luminescence measurements.

Raman spectroscopy was used to collect further structural information and study the influence of rare earths on the glass arrangement. Fig. 7 shows the Raman spectrum of PGBG. As expected for a glassy material, large bands are observed enveloping all contributions of Ge-O, Ga-O, Bi-O and Pb-O bonds. The Raman spectrum is similar to other germano-gallate glasses [START_REF] Zhang | Effects of PbF 2 doping on structure and spectroscopic properties of Ga 2 O 3 -GeO 2 -Bi 2 O 3 -PbO glasses doped with rare earths[END_REF], [START_REF] Yang | Broadband 1.53 μm emission in Er 3+ -doped Ga-Bi-Pb-Ge heavy metal oxide glasses[END_REF], [START_REF] Fukumi | Structures of alkali niobium gallate glasses[END_REF], [START_REF] Mckeown | Raman spectroscopic studies of BaO-Ga 2 O 3 -GeO 2 glasses[END_REF], [START_REF] Montesso | Synthesis and structural characterization of a new SbPO 4 -GeO 2 glass system[END_REF], [START_REF] Subhadra | Infrared and Raman spectroscopic studies of alkali bismuth borate glasses: evidence of mixed alkali effect[END_REF], [START_REF] Saine | Etude vibrationnelle d′aluminates et de gallates de terres rares-III. Aluminates et gallates de structure grenat[END_REF] with predominant absorption bands around 400 and 750 cm -1 . The deconvolution analysis was done based on methods reported elsewhere [START_REF] Montesso | Synthesis and structural characterization of a new SbPO 4 -GeO 2 glass system[END_REF], [START_REF] Miyaji | Structure of PbO-Bi 2 O 3 -Ga 2 O 3 glasses[END_REF], [START_REF] Man | Praseodymium-doped alkali bismuth gallate glasses[END_REF], [START_REF] Guèrineau | Extended germano-gallate fiber drawing domain: from germanates to gallates optical fibers[END_REF] and all band assignments are summarized in Table 2. The bands at 130 cm -1 and 199 cm -1 (group modes) relate to BiO 6 or PbO n units. The signals of each unit were difficult to separate due to the similarities between the GeO 4 and GaO 4 polyhedra, thus both contributions were treated together and called as Ga(Ge) [START_REF] Man | Praseodymium-doped alkali bismuth gallate glasses[END_REF] . The band around 400 cm -1 corresponds to bonds from Ga(Ge)-O-Bi (or Pb), whereas the bands between 525 and 576 cm -1 are associated with vibrations of GaO 4 and GeO 4 polyhedra. At higher frequencies, we identified the stretching modes of Ga(Ge)-O -involving non-bridging oxygen [START_REF] Montesso | Synthesis and structural characterization of a new SbPO 4 -GeO 2 glass system[END_REF], [START_REF] Guèrineau | Extended germano-gallate fiber drawing domain: from germanates to gallates optical fibers[END_REF] .

The inset in Fig. 7 shows the Raman spectrum of the sample containing rare earths. It is evident that the vibrational spectrum could not be explained because the Raman photodetector is much more sensitive to emissions from the rare earths than those arising from the vibrations. Other measurements using a laser with λ = 1064 nm were done to overcome such an inconvenience, but emissions from Tm 3+ ions became significant at this wavelength and no spectra from the vibrational part were observed.

On the other hand, micro-luminescence signals arising from the crystal and glass phases were collected and mapped using the Raman spectrometer, as displayed in Fig. 8. The emission from both phases were in accordance with the elemental mapping results. For the spectrum collected on the glass phase surface, which contained no crystals, the emission curves were typical of amorphous materials. However, when the spectrum was taken from the crystal phase surface, the emission was similar to those arising from crystals, in which the stark levels from Er 3+ ions are more evident [START_REF] Li | Tunable magnetic and fluorescent properties of Tb3Ga5O12 nanoparticles doped with Er3+, Yb3+ and Sc3+[END_REF], [START_REF] Kasprowicz | Up-conversion emission in KGd(WO 4 ) 2 single crystals triply-doped with Er 3+ /Yb 3+ , Tb 3+ /Yb 3+ /Tm 3+ and Pr 3+ /Yb 3+ /Tm 3+ ions[END_REF] . The main emissions were depicted as the energy level diagram in the insert of Fig. 8a. Using the Raman mapping mode, it was possible to generate 3D images of the cubic crystals. As seen in Fig. 8b, the crystal luminescence was zdependent. More importantly, and corroborating with the SEM mapping (Fig. 6), the cubic crystals were confirmed not to be fully organized and the amorphous phase (glass) was present at their center. The emission distribution (from ~650 up to 680 nm) on the cubes shown in Fig. 8 correspond to Er 3+ ions. The emission intensities (indicated by the color scale) decreased as the focal point was shifted towards the crystals, being characterized by green and blue colors. It is worth pointing out that such measurements were performed in a 3D mode, and the resulting curves represent the average data from both phases.

The photoluminescence was also mapped with excitation centered at 1064 nm. Such wavelength corresponds to the tail of the absorption band of Yb 3+ ions in the infrared region, as can be seen in Fig. SM1. The band centered at 980 nm is quite large and appears extended to approximately 800 nm at room temperature. This enables the direct energy transfer from Yb 3+ to Tm 3+ ions, resulting in the emission around 800 nm shown in Fig. 9a. This band arises mainly from 3 H 4 → 3 H 6 transition of Tm 3+ ions [START_REF] Babu | Upconversion luminescence in Tm 3+ /Yb 3+ co-doped lead tungstate tellurite glasses[END_REF] . We could not rule out any contribution from level 4 I 9/2 → 4 I 15/2 transitions of Er 3+ ions, however, the direct emission of Tm is expected to be much more intense than those from Er 3+ [START_REF] Kasprowicz | Up-conversion emission in KGd(WO 4 ) 2 single crystals triply-doped with Er 3+ /Yb 3+ , Tb 3+ /Yb 3+ /Tm 3+ and Pr 3+ /Yb 3+ /Tm 3+ ions[END_REF], [START_REF] Babu | Upconversion luminescence in Tm 3+ /Yb 3+ co-doped lead tungstate tellurite glasses[END_REF] . We also observed an opposite behavior from Fig. 8a, and the most intense signals were observed for glassy phase rather than for the crystals, corroborating with the fact that the rare earth ions were also dispersed within the glassy phase. As already discussed, this work aimed to obtain transparent glass-ceramics containing rare earth-doped active crystals. The relative content of rare earths was set at 2 mol% of Yb 2 O 3 , 0.5 of Er 2 O 3 and 0.5 of Tm 2 O 3 due to the well-known fact that Yb 3+ ions are good IR sensitizers for the upconversion photoluminescence of Er 3+ and Tm 3+ ions [START_REF] Strohhöfer | Absorption and emission spectroscopy in Er 3+ -Yb 3+ doped aluminum oxide waveguides[END_REF], [START_REF] Auzel | Upconversion and anti-stokes processes with f and d ions in solids[END_REF] . To study the upconversion effect, the sample PGBG6 was excited using a solid-state laser at 980 nm and the upconversion emission curves are shown in Fig. 10.

When upconversion takes place emissions can be expected from the excited state of Er 3+ and Tm 3+ . From upper levels of Er 3+ , emission involving a two-photon absorption process result in signals in the green (525 and 550 nm) and red (660 nm) regions, corresponding to 2 H 11/2 → 4 I 15/2 , 4 S 3/2 → 4 I 15/2 and 4 F 9/2 → 4 I 15/2 transitions, respectively. We would also expect two upconversion emissions (at 480 and 800 nm) arising from Tm 3+ ions, as already observed in other glass systems [START_REF] Kasprowicz | Up-conversion emission in KGd(WO 4 ) 2 single crystals triply-doped with Er 3+ /Yb 3+ , Tb 3+ /Yb 3+ /Tm 3+ and Pr 3+ /Yb 3+ /Tm 3+ ions[END_REF], [START_REF] Babu | Upconversion luminescence in Tm 3+ /Yb 3+ co-doped lead tungstate tellurite glasses[END_REF], [START_REF] Ledemi | Multicolor up conversion emission and color tunability in Yb 3+ /Tm 3+ /Ho 3+ triply doped heavy metal oxide glasses[END_REF] , however, in our case, the blue emission corresponding to 1 G 4 → 3 H 6 transition is very weak, but it could still be observed at ~480 nm [START_REF] Babu | Upconversion luminescence in Tm 3+ /Yb 3+ co-doped lead tungstate tellurite glasses[END_REF] .

A three-photon process could also originate an upconversion emission from Er 3+ ions at 410 nm (transition 2 H 9/2 → 4 I 15/2 ), but this emission was not detected in our experiments. The week blue emission of Tm 3+ can be explained by two effects: The first one is related to competition for photons with the Er 3+ ions. The emission at 480 nm corresponding to Tm 3+ is a three-photon dependent process, whereas the upconversion concerning the Er 3+ ions is two-photon dependent process. The second effect is that the absorption tail of the glass matrix falls in the same region where emissions occur, as revealed in Fig. SM1. These two combined effects could explain the low emission intensity of thulium in the blue region. On the other hand, the near infrared emission of Tm 3+ is very intense and observed at ~800 nm ( 1 G 4 → 3 H 5 ) being concomitant with the emission of Er 3+ , as explained above. However, the higher intensity of this band can indicate a direct energy transfer from Yb 3+ to Tm 3+ ions, thus favoring the photoluminescence emission. Fig. 10b shows the well-known Er 3+ emission (downshifting) centered at 1550 nm in the near infrared after excitation at 980 nm, corresponding to the 4 I 13/2 → 4 I 15/2 transition. A broad emission band was observed, corresponding to erbium ions in the amorphous phase, along with sharper emissions around 1535 and 1554 nm ascribed to the stark levels of erbium incorporated in the crystalline phases. These results demonstrate that both upconversion and downshifting processes happen simultaneously and compete for photons, but the green upconversion emission is still high, as can be seen in Fig. 11. The microscope image shows the green emission arising mainly from crystalline phase when the sample is excited at 980 nm with a solid-state laser. This observation corroborates with the statements discussed in this paper that the rare earths were embedded both in the crystalline and glassy phases. The emissions arising from Er 3+ and Tm 3+ ions are of practical interest. While erbium-doped systems have been largely used as probes in nanothermometers (using the upconversion effect) [START_REF] Brites | Thermometry at the nanoscale[END_REF], [START_REF] Cui | Enhanced up-conversion luminescence and optical thermometry characteristics of Er3+/Yb3+ co-doped Sr10(PO4)6O transparent glass-ceramics[END_REF] and telecom systems (downshifting) [33] , Tm 3+ have been studied for white light generation [START_REF] Ledemi | Multicolor up conversion emission and color tunability in Yb 3+ /Tm 3+ /Ho 3+ triply doped heavy metal oxide glasses[END_REF] and bioimaging in the near infrared (750-1000 nm) of biological tissues due to deeper light penetration and higher image contrast [START_REF] Nyk | High contrast in vitro and in vivo photoluminescence bioimaging using near infrared to near infrared up-conversion in Tm 3+ and Yb 3+ doped fluoride nanophosphors[END_REF], [START_REF] Nassar | Niobium oxide doped with Tm 3+ and Gd 3+ ions for multimodal imaging in biology[END_REF] . Concerning the energy level diagram, the mechanisms for energy transfer from Yb 3+ towards both Er 3+ and Tm 3+ ions are well-known [START_REF] Méndez-Ramos | Yb 3+ -Er 3+ -Tm 3+ co-doped nano-glassceramics tuneable up-conversion phosphor[END_REF], [START_REF] Wang | Triple-doped KMnF 3 :Yb 3+ /Er 3+ /Tm 3+ nanocubes: four-color upconversion emissions with strong red and near-infrared bands[END_REF], [START_REF] Mateos | Ultraviolet and visible emissions of Er 3+ in KY(WO 4 ) 2 single crystals co-doped with Yb 3+ ions[END_REF] . Fig. 12 summarizes all emissions that were discussed herein throughout this manuscript.

Conclusions

This study demonstrates the possibility of obtaining Er 3+ /Tm 3+ -doped cubic Yb 3 Ga 5 O 12 single-crystals dispersed in a germano-gallate glass system. Crystals with sizes ranging from 10 to 150 µm were obtained. The REE were present as substitutes of Yb in the crystal lattice. Both glassy and crystal phases exhibited emissions from the RE ions. Upconversion in the visible and near infrared ranges was observed when the glass-ceramic was excited at 980 nm. These materials displayed high sensitivity, thus being promising candidates as optical thermometers and devices for near infrared imaging.

CRediT

Declaration of Competing Interest :

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

Fig. 3 .

 3 Fig. 3. Experimental (blue-PGBG6), simulated (light blue) and reference (PDF-ICDD: 73-1373) (pink-Yb3Ga5O12) diffractograms of PGBG6. Insert: XRD pattern of amorphous PGBG.

Fig. 4 .

 4 Fig. 4. 3D representation of crystal growth obtained by molecular modeling from PDF-ICDD: 73-1373 data using the Mercury software.

Fig. 5 .

 5 Fig. 5. SEM-EDS spectrum of crystal surface.

Fig. 7 .

 7 Fig. 7. Raman spectrum and Gaussian deconvolution of PGBG. Inset: Raman spectrum of PGBG6 corresponding to emissions of rare earth ions under excitation at λ = 632.8 nm.

Fig. 8 .

 8 Fig. 8. a) Micro-luminescence emission spectra (downshifting, λexc= 638 nm) of glass and crystal phases obtained using a Micro-Raman spectrometer and correspondent energy level diagram of Er 3+ ions. b) Luminescence mapping images of crystals in different z positions. The color scale corresponds to emission intensities increasing from blue to red.

Fig. 9 .

 9 Fig. 9. a) Upconversion spectra from glass and crystals under excitation at 1064 nm. The insert shows the energy levels of Tm 3+ ions involved in the emission process. b) Luminescence map obtained at 1064 nm.The color scale corresponds to emission intensities increasing from blue to red.

Fig. 10 .

 10 Fig. 10. a) Upconversion spectra of PGBG6 containing Er 3+ /Tm 3+ -doped gallate crystals under excitation at 980 nm with two different potencies. The insert shows the low intensity upconversion of Tm 3+ at 480 nm. b) Near infrared downshifting emission of Er 3+ ions.

Fig. 11 .

 11 Fig. 11. Microscope images of glass-ceramic sample (PGBG6) before (left image) and during (right image) exposure to 980 nm solid state laser (shown in the center). Green light corresponds to upconversion emissions of Er 3+ ions mainly incorporated in the crystalline phase.

Fig. 12 .

 12 Fig.12. Energy level scheme of Yb 3+ /Er 3+ and Yb 3+ /Tm 3+ under excitation at 980 nm demonstrating both upconversion and downshifting processes.

  authorship contribution statement : Antonio E. Souza: Performed the experiments and wrote the draft. Selma G. Antônio: Performed XRD and simulation studies. Sidney J.L. Ribeiro: Supervised the luminescent measurements and help in the discussion. Douglas F. Franco: Performed Raman measurements and discussion of the data. Gustavo Galeani: Performed luminescence measurements and help in the discussion. Marc Dussauze: Performed micro-luminescence measurements and treated the data and corrected the paper. Thierry Cardinal: Participated in the discussion of the manuscript and corrected the text. Marcelo Nalin: Writing, Editing and Reviewing.

Acknowledgments : The authors acknowledge São Paulo Research Foundation -FAPESP (Brazil) grant numbers 2013/07793-6 and 2019/01223-0 and Conselho Nacional de Desenvolvimento Cientfico e Tecnologico -CNPq (Brazil) (502391/2014-6) for the financial support. Appendix A. Supplementary material