1 Introduction

Rare-earth-doped fiber (REDF) lasers and amplifiers offer numerous advantages, including light weight, compact design, and high electro-optical conversion efficiency. These attributes render them highly suitable for various space applications such as laser communication, Light Detection and Ranging (LIDAR), and space debris removal [1,2,3,4,5]. However, the harsh radiation environment in space can significantly increase optical loss and noise, causing a notable decline in slope efficiency and overall gain of the output laser [6]. These challenges present considerable obstacles to the long-term stability of fiber lasers and amplifiers in space-based operations.

Exposure to high-energy irradiation is known to generate free electron–hole pairs within fiber optic materials. These pairs can be trapped by inherent atomic defects or impurities, forming color centers [6,7,8,9]. The formation of color centers altered the structure and properties of the glass at the microscopic level, particularly influencing the energy level configuration and electronic state distribution of rare earth ions. These changes led to significant modifications in the optical properties of the glass, including the appearance of new absorption and fluorescence bands, a reduction in fluorescence intensity, and increased optical losses due to additional light absorption and scattering [10]. Consequently, these changes substantially affect the transmission and amplification of light in REDFs [11, 12]. Current research on the irradiation resistance of REDFs primarily focuses on understanding the mechanisms underlying radiation damage of fiber materials (in radiation environments), alongside devising effective strategies (via developing diverse radiation-resistant methods) for mitigating such damage, thereby prolonging the service life of optical fibers [10,11,12,13,14]. Notably, two key areas have been emphasized: (1) Elucidating irradiation damage mechanisms, encompassing the evolution of irradiation-induced attenuation (RIA), irradiation-induced emission, and alterations in the internal microstructure of optical fibers [15,16,17]; (2) Defect suppression, wherein the prevailing strategy involves doping with metastable valence ions coupled with a fiber-carrying gas (H2/D2). This approach aims to reduce the formation and propagation of defects in the fiber, thereby enhancing its irradiation resistance [18,19,20]. Nevertheless, although the aforementioned method considerably improves the irradiation resistance of REDFs, a notable concern arises regarding gas leakage. Gas leakage can significantly degrade the irradiation resistance of REDFs, rendering them incapable of meeting the stringent requirements for long-term missions in outer space [21].

The deterioration of optical fiber materials in irradiated environments primarily stems from the accumulation and evolution of irradiation-induced point defects, leading to a decline in laser performance [22]. Thus, boosting the compounding efficiency of irradiation-induced point defects (such as interstitial atoms and vacancies) can significantly bolster the irradiation resistance of optical fibers [22,23,24,25]. Relevant studies have demonstrated that introducing second-phase particles into a material (with a specific density and uniform size), to create phase interfaces with higher density is a crucial strategy for enhancing the composite efficiency of irradiation-induced point defects [26,27,28]. These phase-interface regions can serve as composite centers for irradiation-induced point defects by absorbing radiant energy and trapping damaged defects, thereby mitigating radiation-induced damage [29,30,31]. According to the Mie scattering theory, when the size of second-phase particles remains below 100 nm, light-scattering losses are negligible. Hence, incorporating second-phase particles, characterized by a specific density and a homogeneous size below 100 nm in diameter, into the core glass is anticipated to be a novel approach for enhancing the irradiation resistance of REDFs [32].

In this study, erbium-ytterbium co-doped high-phosphorus silica glasses with nanoscale phase-separated structures were synthesized using a sol–gel method combined with phase-separation engineering. By controlling the annealing duration, we modulated the redevelopment of the phosphorus-rich phase in the glass to form nanoscale phase interfaces with different distribution densities between the phosphorus-rich and silica-rich phases. Using techniques such as high-resolution transmission electron microscopy (HRTEM), nuclear magnetic resonance (NMR), and spectral characterization, we tracked the evolution of internal microstructures of the glasses at the atomic level. Our findings confirmed that the Si-O-Si bonds in the glass matrix underwent partial cleavage as annealing progressed, thus resulting in their conversion to P-O-Si and P-O-P bonds. This conversion directly results in increased density of the phase interface inside the glasses. Moreover, after X-ray irradiation at a total dose of 1 kGy, the RIA and photoluminescence intensity of the annealed glass improved significantly compared to those of pristine glass. This promotes the effectiveness of increasing the phase-interface density for enhancing the radiation resistance of glasses. Finally, we conducted online irradiation experiments utilizing the MEY-series of fibers developed in our laboratory. The term “MEY-series fibers” specifically refers to EYDFs prepared using the Modified Chemical Vapor Deposition (MCVD) method. The results revealed that the RIA at 1080 nm was reduced by 45% for MEY3 (annealed for 3 h) and 73% for MEY20 fibers (annealed for 20 h) with different phase-interface densities compared to pristine fibers. This study suggests a novel method to enhance to radiation-resistance behaviors of REDFs.

2 Experimental details

2.1 Sample preparation

Erbium-ytterbium-co-doped high-phosphorus silica glasses were prepared using an improved sol–gel method combined with a high-temperature sintering technique. The experimental flowchart for glass preparation was illustrated in Fig. 1a. For more comprehensive details on the preparation process, please refer to previous studies [33, 34]. TEOS (tetraethyl orthosilicate), C2H5OH, ErCl3·6H2O, YbCl3·6H2O, and H3PO4 were used as precursors, with deionized water facilitating hydrolysis. After sealing and stirring for 20 h at room temperature, a colorless and transparent silica sol was obtained. Subsequently, the silica sol was dried at 80 °C for 10 h, producing transparent dry gel particles. These particles underwent heat treatment in an oxygen environment at 1000 °C for 5 h to decompose hydroxyl groups and organic components. After treatment, the gel particles exhibited irregular shapes and sizes up to several millimeters, resulting in macroscopic defects such as bubbles and inclusions during high-temperature sintering. To improve particle size distribution and morphology, the gel particles were milled using a planetary ball mill to achieve sizes below 15 μm before powder sintering. Finally, the powder was melted at 1700 °C in a vacuum furnace for 1 h, resulting in the production of high-quality optical glass samples. The theoretical composition of the glasses was intended to be 0.04Er2O3-0.44Yb2O3-10P2O5-89.52SiO2 (mol%). However, the actual composition determined via inductively-coupled plasma atomic emission spectrometry (Thermo ICAP 6300) was 0.04Er2O3-0.41Yb2O3-8.15P2O5-91.43SiO2 (mol%), attributable to the volatile nature of elemental phosphorus. The glass samples, comprising these components, were classified into three categories based on variations in the annealing duration: SEY0 glass (pristine), SEY50 glass (annealed for 50 h), and SEY100 glass (annealed for 100 h). All annealing treatments were conducted in a tube furnace under argon protection at an annealing temperature of 850 °C, as elaborated further later herein. The SEY0 and SEY50 glasses underwent X-ray irradiation at doses up to 1 kGy from an X-ray irradiator (X-RAD 160) at a rate of 10 Gy/min in a normal environment. After irradiation, these glasses were denoted as SEY0-1kGy and SEY50-1kGy glasses, respectively.

Fig. 1
figure 1

Experimental flow chart for the preparation of a SEY-series glasses and b MEY-series fibers

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The self-developed MEY0 fiber is a radiation-resistant reactive silica fiber. Its preform was fabricated by combining MCVD and nano-sol doping. Its chemical composition is similar to that of SEY0 glass, with the exception of an additional 8000 ppm of Ce2O3. The preparation process for MEY-series fibers was depicted in Fig. 1b. The fiber coating was removed before annealing. The fibers were defined based on annealing time as follows: MEY0 fibers (unannealed), MEY3 fibers (annealed for 3 h), and MEY20 fibers (annealed for 20 h). Additionally, the fibers were annealed in an argon-protected tube furnace at the same annealing temperature as the glass. The unique rigid glass network of silica-based glass and the annealing temperature well below Tg ensured that there was no change in the valence state of Ce.

Herein, the glass samples SEY0, SEY50, SEY100, SEY0-1kGy, and SEY50-1kGy are collectively referred to as the SEY-series glasses. Similarly, the optical fibers comprising the MEY0, MEY3, and MEY20 fiber samples are denoted as the MEY-series fibers.

2.2 Analyses of samples

The glass transition temperature (Tg) of the unannealed SEY0 glass was ~930 °C, as measured using a differential scanning calorimetry analyzer (DSC, Netzsch STA449F3) at a heating rate of 15 K/min. An HRTEM equipment (Tecnai G2, FEI) equipped with energy-dispersive X-ray spectroscopy (EDS) was employed to evaluate the morphology and size distribution of the nanocrystals in the samples. The HRTEM imaging detector targeted the backscattered electrons, and the counting time resulting from the EDS was ~180 s. Raman spectra were recorded via a Raman spectrometer (Horiba) with a 633 nm He–Ne laser. For structural analysis, an X ‘Pert PRO (Netherlands) diffractometer was utilized to analyze the glass. All the NMR experiments were conducted using a Bruker spectrometer (Avance III HD 500 MHz, 11.7 T). Specifically, 31P Magic Angle Spinning (MAS) NMR experiments were performed at 202.44 MHz, utilizing a 4-mm probe operating at a rotor spinning rate of 12 kHz, a 90° pulse of 2.5 μs, and a relaxation delay of 40 s. Additionally, 31P constant-time (CT) double quantum (DQ)-based dipolar recoupling effects nuclear alignment reduction (DRENAR) experiments were carried out using a 2.5 mm probe and a rotor frequency of 25 kHz. A 90° pulse length of 4.3 μs and a relaxation delay of 10 s were set with a loop value of 2. The absorption spectra of the glasses were recorded using a Lambda 950 ultraviolet (UV)–visible–near-infrared spectrophotometer (Perkin Elmer). Furthermore, an FLS 920 spectrofluorometer (Edinburgh) was employed to detect the photoluminescence spectra and fluorescence decay curves of Yb3+ and Er3+ ions.

The online test was performed using an X-ray machine with a photon energy ranging from 40–100 keV. To account for the potential impact of X-rays on the measuring instrument, we performed an on-line test to evaluate the radiation-induced loss at 1080 nm of the MEY series fibers. This test was conducted under identical conditions using the setup depicted in Fig. S1. Throughout the on-line irradiation test, only the MEY-series fibers were exposed to the radiation field. The length of the measured MEY-series fiber was 30 cm, with the signal power at 1080 nm being 30 mW.

3 Results and discussion

3.1 Microstructural study of glass

Figure 2a–c presents transmission electron microscopy (TEM)/HRTEM images of SEY0 glass, illustrating nanoscale second-phase particles dispersed in a droplet-like morphology throughout the glass matrix. The diameter of the second-phase particles varies from about 3–10 nm. The selected-area electron diffraction data for the second-phase particles in Fig. 2c indicate that they are in an amorphous state (Fig. 2c inset). Figure 2d shows a high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) image and the corresponding elemental mapping image of SEY0 glass. These images demonstrate the selective aggregation of P and Yb elements. Line-scan analysis of SEY0 glass via scanning transmission electron microscopy–energy-dispersive X-ray spectroscopy (STEM-EDS) (see Fig. 2e–g) further confirmed that the main chemical composition of the second-phase particles was an aggregation of rare-earth (Er and Yb) and P elements. Notably, due to the low doping content of Er, which was below the detection limit of the detectors in both HAADF-STEM and STEM-EDS, its presence could not be effectively detected. However, this did not affect the conclusion, as Er and Yb possess similar chemical properties pertaining to their atomic structures. The primary reason for phase separation in SEY0 glasses has been described in detail in our previous study [33]. This phenomenon arises due to the presence of two different types of network formers in the glass: P5+ and Si4+. P5+ exerts a stronger field strength than Si4+, thus rendering it more robust against non-bridging oxygen. Consequently, structural disparities emerge between [SiO4] and [PO4], rendering them incompatible. This incompatibility leads to phase separation, resulting in the formation of regions comprising phosphate and silicate phases.

Fig. 2
figure 2

TEM/HRTEM bright-field images of SEY0 glass at different scales: a 50 nm, b 20 nm, and c 2 nm, with the inset showing an SAED image. d HAADF-STEM image of SEY0 glass accompanied by corresponding elemental mapping (Si, Yb, and P elements). e–g TEM dark-field image of SEY0 glass illustrating the position of the EDS line scan, alongside the scan results of a representative second-phase particle. The glass was ion-thinned prior to TEM analysis

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Figure 3a depicts the DSC curve of SEY0 glass, revealing a Tg of ~930 °C. Typically, for quartz glass with an OH content lower than 1 × 10−6 cm−1, the Tg is ~1207 °C [35]. To further regulate the phase transition behavior of SEY0 glass and facilitate the formation of ordered structures, we selected 850 °C as the annealing temperature. This lower annealing temperature decelerates the molecular motion of the material, allowing for the introduction of ordered structures without compromising the existing glassy state. Consequently, controlled phase-separation kinetics are enabled, promoting the development of desired structural configurations.

Fig. 3
figure 3

a DSC curve of SEY0 glass. TEM/HRTEM bright-field images of SEY50 glass at different scales: b 50 nm, c 2 nm. d Enlarged view of the red-highlighted area shown in this figure

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The size of the second-phase particles in SEY50 glass was notably more pronounced than those in SEY0 glass (~10–20 nm in diameter), even forming irregular lattice fringes in localized regions (see Fig. 3d). This observation suggests that the SEY50 glass possesses a higher phase-interface density. This heightened density can be attributed to the temperature-induced phase-transition behavior of SEY50 glass, which exacerbates the significant migration and reorganization of second-phase particles (phosphorus phase) within the glass. Additionally, the formation of irregular lattice fringes is indicative of changes in the chemical nature or energy state of P. In various instances, minute amounts of P2O5 serve as nucleating agents [36]. Therefore, it can be reasonably inferred that SEY100 glass contains more irregular lattice fringes and higher phase-interface densities than SEY50 glass after prolonged annealing times.

Figure 4a displays the Raman spectra of SEY0, SEY50, and SEY100 glasses, revealing prominent vibrational bands at 430, 480, 615, 720, 808, 1170, and 1320 cm−1 [37]. Among these, the bands at 480 and 615 cm−1 are attributed to Si-O-Si planar quaternary rings and planar ternary rings, respectively [37, 38]. The bands at 430 and 808 cm−1 correspond to the bending and symmetric stretching vibrations of the Si-O-Si bond, respectively [37]. Furthermore, the band near 720 cm−1 is associated with the bending vibration of O=P-O, while vibrations near 1170 cm−1 are typically attributed to Si-O-Si and P-O-Si bond vibrations [38, 39]. The vibrations at 1320 cm−1 mainly arise from the P=O double bond [39].

Fig. 4
figure 4

a Normalized Raman spectra, b XRD patterns, and c full width at half maximum of broad peaks in XRD patterns of SEY0, SEY50, and SEY100 glasses. d, e 31P MAS NMR spectra (yellow) and 1D-refocused INADEQUATE spectra (purple) of SEY0 and SEY100 glasses. f CT-DRENAR-POST-C7 data curves for SEY0 and SEY100 glasses, where solid curves represent the corresponding effective double-spin simulations

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Prolonged annealing significantly altered both the vibrational spectra and X-ray diffraction (XRD) patterns of the glass samples. In particular, the Raman vibrational band at 615 cm−1 displayed gradual attenuation, while the band at 720 cm−1 intensified, indicating considerable changes in the internal atomic structure resulting from elevated annealing temperatures. This change resulted in the development of second-phase particles and enhancement of phase-interface densities. The decrease in intensity at 615 cm−1 might be caused by the partial breakage or reorganization of Si-O-Si bonds within the glass matrix, resulting in diminished vibrational intensity of the Si-O-Si planar ternary rings. Conversely, the increase in strength at 720 cm−1 implies a rise in the prevalence of P-O-P bonding, potentially caused by the migration and reorganization of second-phase particles (phosphorus phase) during annealing. Furthermore, XRD analysis revealed that extended annealing substantially broadened the characteristic peak associated with the amorphous structure (see Fig. 4b, c). This broadening signifies increased short-range order within the amorphous structure, ascribed to the migration and recombination of the secondary-phase particles. Additionally, augmented phase segregation within the glass matrix contributed to the broadening of the XRD peaks. These changes indicate the complex structural transformations occurring in the glasses during the annealing process, highlighting the dynamic nature of their atomic arrangement and phase distribution. Ultimately, these processes contribute to the strengthening of phase-interface densities within the glasses.

Figure 4d, e depicts the 31P MAS NMR spectra (yellow) and one-dimensional (1D)-refocused INADEQUATE spectra (purple) of the SEY0 and SEY100 glasses, respectively. The 1D-refocused INADEQUATE experiment facilitates a straightforward differentiation between phosphate species involved in P-O-P bonding and those lacking such bonds [40, 41]. As illustrated in Fig. 4d, e, the 31P MAS NMR spectra of SEY0 and SEY100 glasses closely resemble their 1D-refocused INADEQUATE spectra, suggesting that among all P atoms within these glasses, at least one P-O-P bond exists, while P(0) units are absent in both cases [42].

Figure 4f presents the CT-DRENAR-POST-C7 decay curves for SEY0 and SEY100 glasses under identical conditions. This experiment can effectively detect the intensity of dipole–dipole coupling between 31P and 31P in glass [43]. To characterize the average strength of the homonuclear dipole–dipole coupling between nucleus \(k\) and the observed spin \(j\), the sum of squares of the experimental dipolar coupling constant, \(\sum _{j}{b}_{{jk}}^{2}\), can be obtained using Eq. 1 [44].

$$\frac{({S}_{0}-S^{\prime} )}{{S}_{0}}=\frac{6}{5}{\left\{\sqrt{\left(\sum _{j}{b}_{{jk}}^{2}\right)N{T}_{r}}\right\}}^{2}\left(1+\cos 2\theta \right),$$
(1)

where \(k\) represents all the spins coupled with the observed spin \(j\); \({b}_{{jk}}\) is the dipolar coupling constant between \(j\) and \(k\), separated by distance \({r}_{{jk}}\); and \(\theta\) denotes the phase rotation angle (0° ≤ 2\(\theta\) ≤ 360°), increased in discrete steps of 18°. In Fig. 4f, the value of \(\sum _{j}{b}_{{jk}}^{2}\) is ~25.6 × 105 Hz for SEY100 glass, while it is only ~8.1 × 105 Hz for SEY0 glass. According to literature, the value of \(\sum _{j}{b}_{{jk}}^{2}\) for AlPO4 glass is ~3.3 × 105 Hz, whereas the values for both SEY0 and SEY100 glasses exceed 3.3 × 105 Hz [39]. This indicates the presence of P-O-P bonds in both glasses. However, the number of P-O-P bonds in SEY0 glass is less than one-third of that in the SEY100 glass. Such a significant disparity implies a distinct phase-separated structure in SEY100 glass, signifying a tendency toward micro-crystallization. This conclusion is further supported by the pronounced sharpness observed in the 31P MAS NMR spectrum of the SEY100 glass.

3.2 Spectroscopic properties of glass

Figure 5a displays the UV–visible transmittance spectra (200–800 nm) of the SEY0, SEY50, and SEY100 glasses. Compared with SEY0 glass, both SEY50 and SEY100 glasses exhibited considerable decreases in the UV–visible transmittance, primarily due to the increased size of the second-phase particles (refer to Figs. 2a and 3b). Additionally, compared to the SEY0 glass, a noticeable blue shift was observed in the transmittance spectra of SEY50 and SEY100 glasses at the UV cutoff edge. This shift primarily stems from alterations in the internal structure of the glass, resulting in changes in the energy band structure. Figure 5b illustrates the absorption spectra (850–1650 nm) of the SEY0, SEY50, and SEY100 glasses. The overall intensities of the absorption spectra of SEY50 and SEY100 glasses were significantly lower than those of the SEY0 glass. However, the overall intensities of the absorption spectra of SEY50 and SEY100 glasses were almost identical. Figure 5c, d depict the fluorescence spectra (850–1650 nm) and normalized fluorescence spectra (used to measure the energy-transfer efficiency from Yb3+ to Er3+) of the SEY0, SEY50, and SEY100 glasses pumped using an 896 nm Xenon lamp. Typically, the energy-transfer efficiency between Yb3+ and Er3+ can be effectively determined from the relative fluorescence intensities of Yb3+ and Er3+ ions [45, 46]. With the prolongation of the annealing time, the fluorescence spectra of Yb3+ and Er3+ both showcased a decreasing trend in the overall intensity (see Fig. 5c). However, the energy-transfer efficiency from Yb3+ to Er3+ increased (see Fig. 5d). This is mainly due to the relatively weak fluorescence intensity of SEY50 and SEY100 glasses at 975 nm compared with SEY0 glass. The growth of second-phase particles induces more rare-earth ions to aggregate inside them, considerably reducing the distance between Yb3+ ions and Er3+ ions. Consequently, this enhances the energy-transfer efficiency between Yb3+ and Er3+ ions [33]. Meanwhile, the abundant P=O double bonds in the second-phase particles have high phonon energies. This is conducive for suppressing the reverse energy transfer from the excited state of Er3+ to Yb3+ ions, further improving the energy-transfer efficiency from Yb3+ to Er3+ ions.

Fig. 5
figure 5

a UV–visible transmission spectra (200–800 nm), b Absorption spectra (850–1650 nm), c Fluorescence spectra (excitation at 896 nm), and d Normalized fluorescence spectra of SEY0, SEY50, and SEY100 glasses. The test thickness of all glass samples was 2 mm

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Table 1 presents the fluorescence lifetimes of Yb3+ ions (at 1030 nm) and Er3+ ions (at 1.5 μm) in SEY0, SEY50, and SEY100 glasses. The decay behavior of numerous materials deviates from the exponential decay model [47, 48]. Therefore, for such materials, a common approach for fitting decay is to leverage multi-exponential fitting, followed by calculating the average decay lifetime based on this approach. Two prevalent methods for calculating the average lifetime are the amplitude-weighted average lifetime and intensity-weighted average lifetime [49, 50]. The intensity-weighted average lifetime emphasizes the influence of long-lived components, thereby mitigating the impact of short-lived components on the fitting results. This ensures greater stability in the quantity changes of fluorescence components throughout the fitting process. Hence, this method finds widespread application in perovskites, quantum dots, and nanocrystalline glasses [50,51,52]. The specific calculation formula is expressed in Eq. (2).

$$\tau =\frac{{\sum }_{i}^{n}{A}_{i}{\tau }_{i}^{2}}{{\sum }_{i}^{n}{A}_{i}{\tau }_{i}}$$
(2)

Here, \({\tau }_{i}\) represents the component of each lifetime in the intensity-weighted average lifetime, and \({A}_{i}\) is the fractional intensity corresponding to that lifetime component. In erbium-ytterbium co-doped high-phosphorus silica glass, characteristic phase separation may lead to the simultaneous presence of multiple luminescent groups, including phosphorus and silicon phases. However, because the fluorescence lifetimes of Er3+ in both the phosphorus and silicon phases are on the order of ~10 ms, they can be fitted directly using a single index [53, 54]. The drastic decrease in the fluorescence lifetimes of Yb3+ and Er3+ ions indicates that the excessively annealed SEY100 glass exhibited the typical phenomenon of quenching in rare-earth-ion concentrations. This is because excessive heat treatment results in the inordinate aggregation of rare-earth ions within the second-phase particles, causing a “blocking” phenomenon, which in turn triggers the quenching of rare-earth ions. Due to the severe degradation of the fluorescence lifetime performance of SEY100 glass, we performed irradiation resistance experiments using only SEY0 and SEY50 glasses.

Table 1 Fitting results of fluorescence decay curves for Yb3+ and Er3+ in SEY-series glasses
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3.3 Evaluation of radiation resistance in fibers

To validate the effectiveness of the phase-boundary strategy in enhancing the radiation resistance of the glasses, SEY0 and SEY50 glasses were selected for X-ray irradiation at a total dose of ~1 kGy. Figure 6a, b shows the RIA and fluorescence spectra (excitation at 980 nm) for the SEY0-1kGy and SEY50-1kGy glasses, respectively. RIA is a critical parameter for assessing the performance of radiation-resistant materials. It can be derived by subtracting the absorption coefficient of an irradiated sample from that of a pristine sample. According to the pertinent literature, an increase in RIA in REDFs is associated with a change in the valence state of rare-earth ions and the formation of matrix-associated point defects [55, 56]. Notably, point defects are recognized as the cause of radiation-induced darkening in REDFs. The RIA intensity in the UV-visible region of SEY0-1kGy glass was higher than that of SEY50-1kGy glass, indicating a greater irradiation attrition coefficient for the former. However, this difference in RIA intensity was only observed in 2 mm thick glass samples. For optical fibers of significant length, the cumulative effect of RIA would have been substantial. Additionally, the reduction in the photoluminescence intensity of SEY50-1kGy glass was considerably less pronounced compared to that of SEY0-1kGy glass. Furthermore, the decline of Yb3+ and Er3+ fluorescence lifetimes in SEY50-1kGy glass was weak compared to that observed in SEY0-1kGy glass (see Table 1). These findings suggest that SEY50 glass exhibits higher radiation resistance than SEY0 glass.

Fig. 6
figure 6

a RIA spectra and b Fluorescence spectra (excitation at 980 nm) of SEY0-1kGy and SEY50-1kGy glasses. c Online irradiation setup for MEY-series fibers. d Radiation-induced attenuation vs. radiation dose at 1080 nm for MEY-series fibers under online irradiation

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The effect of the phase-interface strategy on the radiation resistance of Er-Yb co-doped high-phosphorus silica fibers was further investigated. We prepared an MEY-series of fibers with different phase-interface densities by adjusting the annealing time of the fibers. This was achieved by removing the coating on the MEY fibers prior to annealing and then allowing them to anneal in an argon atmosphere. Depending on the annealing time, they were denoted as MEY0 fibers (pristine), MEY3 fibers (annealed for 3 h), or MEY20 fibers (annealed for 20 h). Using the setup in Fig. 6c, we performed an online test of the radiation-induced attenuation of the MEY-series fibers at 1080 nm as a function of the X-ray dose under the same conditions. Throughout the process, only the MEY-series fibers were exposed to the radiation field. The length of the measured MEY-series fibers was 30 cm, with the signal power at 1080 nm being 30 mW. The results of the online irradiation testing are shown in Fig. 6d. At a dose rate of 100 rad/min and a cumulative radiation dose of 10 krad, the RIA at 1080 nm for the MEY0, MEY3, and MEY20 fibers were ~1.1, 0.6, and 0.3 dB, respectively. Compared to the MEY0 fibers, the induced attenuation of the MEY3 and MEY20 fibers decreased by 45 and 73%, respectively. This indicates that introducing an interfacial phase can effectively enhance the radiation resistance characteristics of the MEY fibers.

4 Conclusion

In summary, we prepared erbium-ytterbium co-doped high-phosphorus silica glass with nanoscale phase-separated structures based on the sol–gel method combined with phase-separation engineering. The emergence of phase-separated structures in the glass matrix results in the coexistence of nanoscale phosphorus-rich and silica-rich phases, leading to the formation of homogeneously distributed nanoscale phase interfaces. The density of the phase interface inside the glass was enhanced by varying the annealing time. Following X-ray irradiation at a total dose of 1 kGy, the RIA and spectral properties of pristine and annealed glasses were compared. The results confirmed that enhancing the phase-interface density significantly increased the radiation resistance of the glasses. To further investigate the applicability of the conclusion above to erbium-ytterbium co-doped high-phosphorus silica fibers, we conducted online irradiation tests on MEY-series fibers with different annealing treatments, that is, with different phase-interface densities. The results confirmed that compared with pristine fibers, the RIA at 1080 nm reduced by 45% and 73% for the MEY fibers annealed for 3 h and 20 h, respectively. These findings provide an innovative method to enhance the radiation resistance of rare earth doped fibers, presenting a new scheme for the development of novel radiation-resistant rare earth doped fibers.