On-chip integrated multicolor microlasers have been intensively studied owing to their practical applications ranging from miniaturized coherent light source, optical sensing, to visible light communication [1–7]. Compared with the combination of several microlasers, the white-light microlaser from a single nanomaterial is more attractive for on-chip integration [8–16]. In 2015, Fan et al. reported a monolithic white-light laser by parallelly combining ZnCdSSe-based segments into the same microcrystal [8]. Soon after, multicolor microlasers were widely observed in various systems, e.g., ${\mathrm{V}}_2\mathrm{C}$-based quantum dots, luminescent polymers, organic dyes-doped polymer, and liquid crystals [9–14]. Despite the continuous success, these white-light microlasers are either simple compositions of parallel microlasers or generated by the reabsorption and the re-emission. From the point view of practical applications, these multicolor microlasers usually come from different locations and the eventual white colors are strongly dependent on the excitation density.

Benefiting from their intrinsic features of plentiful energy levels and long decay lifetimes, lanthanide (${\mathrm{Ln}}^{3+}$) ions have shown their ability of photon emission across the entire ultraviolet-visible (UV-vis) spectrum [17–22]. Two main approaches have been proposed to utilize the characteristics of ${\mathrm{Ln}}^{3+}$ in on-chip lasing actions [23–29]. The first one is the combination of ${\mathrm{Ln}}^{3+}$ ions with on-chip integrated microcavities with ultrahigh quality ($\mathit{Q}$) factors, e.g., microtoroids [23–29]. By preserving $\mathit{Q}$ factors at $\sim {10}^7$, on-chip integrated microlasers have been successfully demonstrated at near-infrared, green, and blue light wavelengths [23–27]. Nevertheless, the lasing performance at the short wavelength is highly constrained by its low doping concentration, small absorption cross section, and large nonradiative relaxation loss due to the direct interaction between ${\mathrm{Ln}}^{3+}$ ions and silica matrix [23,24]. Recently, the core-shell structures [22,30–33] have been developed to tackle these problems and improve the lasing actions [17,34–36]. However, the size of upconversion nanocrystals (UCNCs) is relatively large, introducing strong scattering loss and constraining the intrinsic $\mathit{Q}$ factors of ${\mathrm{Ln}}^{3+}$-based microresonators. Therefore, the development of stable white-light lasers is facing a severe challenge of the trade-off between high doping concentration and high-$\mathit{Q}$ microcavity. Herein, by the introduction of cunning ${\mathrm{CaF}}_2$ NCs and the controlled growth of ${\mathrm{Ln}}^{3+}$-doped NCs, we experimentally demonstrate the high-$\mathit{Q}$ UCNCs-doped microtoroid. The high-$\mathit{Q}$ factors improve the blue, green, and red lasers to similar thresholds, resulting in stable white-light microlasers.

The controlled growth of ${\mathrm{CaF}}_2\text{:}{\mathrm{Yb}}_{35}{\mathrm{Tm}}_{1.5}{\mathrm{Er}}_{0.5}$ UCNCs in the silica film was implemented through a solgel technique followed by a spin-coating operation (see Section 4) [37,38]. The doping concentration of ${\mathrm{Ln}}^{3+}$ ions is deduced from the stoichiometric amounts of precursors in the fabrication process. Note that the efficient ${\mathrm{CaF}}_2$ host [39,40] and CMOS-compatible ${\mathrm{SiO}}_2$ matrix [41–44] were chosen for their similar group refractive index values. Figure 1 shows the structural characterization of UCNCs-doped silica glass film with a thickness of around 630 nm. This nanofilm obtained from the optimized growth conditions exhibits an average variation in the vertical direction of less than 20 nm. The transmission electron microscopy (TEM) image in Fig. 1(a) reveals that NCs with an average size of 4.3 nm were uniformly separated in the silica glass matrix. The high-resolution TEM image [HRTEM, Fig. 1(b)] and the corresponding fast Fourier transforms [FFTs, Fig. 1(c)] confirm the single-crystalline structure of NCs, demonstrating that ${\mathrm{CaF}}_2$ crystallized with (111) preferred orientation in the silica glass matrix [39,40,45]. This is consistent with the X-ray diffraction (XRD) data in Fig. 1(e). Indeed, four diffraction peaks due to ${\mathrm{CaF}}_2$ (JCPDS card #35-0816) with a pure cubic phase can be observed from the XRD pattern except for a broad band centered at 22°, which can be ascribed to amorphous ${\mathrm{SiO}}_2$. Moreover, there are no obvious cracks and roughness even from the scanning electron microscopy [SEM, Fig. 1(d)] result. Such ultrasmooth surface of glass film would determine the $\mathit{Q}$ factor of the resultant UCNCs-doped microtoroid cavity [46].

The chemical composition of the glass film can be verified according to the energy-dispersive X-ray spectroscopy (EDS), UV-vis, and photoluminescence (PL) results. Only Si, O, Ca, and Yb elements can be detected from the EDS measurement [Fig. 2(a)]; however, there are characteristic emitting peaks associated with ${\mathrm{Tm}}^{3+}$ and ${\mathrm{Er}}^{3+}$ ions from the PL spectra later shown in Fig. 5(b). In addition, two broad absorption peaks [Fig. 2(c)] were observed at around 973 and 1520 nm, which can be attributed to the ${{}^2\mathrm{F}}_{7/2}\to {}^2{\mathrm{F}}_{5/2}$ transition of ${\mathrm{Yb}}^{3+}$ ions, and the ${}^4{\mathrm{I}}_{15/2}\to {}^4{\mathrm{I}}_{13/2}$ transition of ${\mathrm{Er}}^{3+}$ ions, respectively. No second phase can be identified from the XRD pattern of the glass film after doping ${\mathrm{Ln}}^{3+}$ (i.e., ${\mathrm{Yb}}^{3+}$, ${\mathrm{Tm}}^{3+}$, and ${\mathrm{Er}}^{3+}$) ions. This clearly suggests the successful incorporation of ${\mathrm{Ln}}^{3+}$ ions into the ${\mathrm{CaF}}_2$ host lattice.

To understand the crystallization of the glass matrix, we carried out thermogravimetry/differential thermal analysis (TG/DTA) and XRD measurements. The amorphous-to-crystalline phase transition was also consistently recorded in TEM results (Fig. 3). Together with the exothermic/endothermic peaks in the DT profile, the TG curve of xerogel in Fig. 2(d) shows four stages of weight loss in air, which can be ascribed to the evaporation of water and organic solvent between 25°C and 110°C, the burnout of excess acetic acid and ethanol between 110°C and 260°C, the complete combustion of polyester via the hydrolysis reaction of tetraethyl orthosilicate (TEOS) and acetates between 260°C and 400°C, and the crystallization of the cubic ${\mathrm{CaF}}_2$ phase between 400°C and 800°C, respectively. As evidenced by the XRD patterns [Fig. 2(e)] and TEM images (Fig. 3) for a series of ${\mathrm{CaF}}_2\text{:}{\mathrm{Yb}}_{35}{\mathrm{Tm}}_{1.5}{\mathrm{Er}}_{0.5}\text{-}{\mathrm{SiO}}_2$ nanofilms annealed at different temperatures, the films remain amorphous below 400°C accompanying with the phase separation and nucleation of amorphous clustering ions at this stage, which satisfies the nonclassical nucleation theory [47]. Then, well-organized crystalline grains can be obtained at a relatively high temperature (i.e., 800°C). This clearly suggests the occurrence of crystallization and growth of NCs, which follows the CN theory [47]. From Fig. 2(e), the corresponding diffraction peaks emerging from the broadband spectrum can be ascribed to (111), (200), (220), and (311) planes of the cubic ${\mathrm{CaF}}_2$ NCs, respectively [39,40,45]. Moreover, the impurity phases are observed from the film annealed at 1000°C. Meanwhile, the observed NCs experience obvious increases in size (Fig. 3). These are because heat treatment at elevated temperature promotes the relocation of dopant ions across the initial crystalline surface, thus resulting in the reconstruction and recrystallization of the crystalline grains [47,48].

Accordingly, we have explored the dependence of surface quality on the heating rates during the multistep calcination. On the basis of the TG/DTA results, a group of ${\mathrm{CaF}}_2\text{:}{\mathrm{Yb}}_{35}{\mathrm{Tm}}_{1.5}{\mathrm{Er}}_{0.5}\text{-}{\mathrm{SiO}}_2$ film samples were obtained under different heat treatment before cooling down to room temperature. The photographs of annealed coatings (Fig. 4) visually reflect the evolution of the surface roughness. Clearly, the NCs-doped glass film treated by an optimized four-stage calcination [Fig. 4(d)] shares high transparency in the visible region [Fig. 2(b)] and ultrasmooth surface finish [Fig. 1(d)]. This is reasonable since the rapid rate, especially at the first heating stage, results in sharp strains from the evaporation, highly restricting the development of ultrasmooth surface. Also, it is worth noting that the sizes, density, and distribution of ${\mathrm{CaF}}_2\text{:}{\mathrm{Ln}}^{3+}$ NCs in glass matrix can be optimized through the precise control of the dopant concentration in the precursor solution, the duration time, and temperature of calcination at each stage.

An extensive optical characterization was performed to validate white-light emission behavior of ${\mathrm{CaF}}_2\text{:}{\mathrm{Ln}}^{3+}\text{-}{\mathrm{SiO}}_2$ glass films, including multicolor emission spanning the visible spectrum, particularly red, green, and blue (RGB) colors. In this study, ${\mathrm{Yb}}^{3+}$ ions are selected as the sensitizer to absorb 980 nm irradiation and then transfer to the neighboring activator ions, while ${\mathrm{Er}}^{3+}$ and ${\mathrm{Tm}}^{3+}$ ions act as the activator to produce RGB emission. Figure 5(b) shows the upconversion multicolor emitting spectra of ${\mathrm{CaF}}_2\text{:}{\mathrm{Yb}}_{35}{\mathrm{Tm}}_{1.5}$, ${\mathrm{CaF}}_2\text{:}{\mathrm{Yb}}_{35}{\mathrm{Er}}_{1.0}$, and ${\mathrm{CaF}}_2\text{:}{\mathrm{Yb}}_{35}{\mathrm{Tm}}_{1.5}{\mathrm{Er}}_{0.5}$ NCs-doped films under the excitation of a continuous-wave (CW) 980 nm laser. The simplified energy level diagram for the observed light is depicted in Fig. 5(a). Mainly, the emission lines can be assigned to blue (${}^1{\mathrm{D}}_2\to {}^3{\mathrm{F}}_4$ and ${}^1{\mathrm{G}}_4\to {}^3{\mathrm{H}}_6$ of ${\mathrm{Tm}}^{3+}$), green (${}^2{\mathrm{H}}_{11/2}\to {}^4{\mathrm{I}}_{15/2}$ and ${}^4{\mathrm{S}}_{3/2}\to {}^4{\mathrm{I}}_{15/2}$ of ${\mathrm{Er}}^{3+}$), and red (${}^1{\mathrm{G}}_4\to {}^3{\mathrm{F}}_4$ of ${\mathrm{Tm}}^{3+}$, and ${}^4{\mathrm{F}}_{9/2}\to {}^4{\mathrm{I}}_{15/2}$ of ${\mathrm{Er}}^{3+}$) bands, respectively [15,16,19,20]. The upconversion mechanism responsible for these characteristic emissions is attributed to energy transfer upconversion processes [15–22]. Emission peaks at the red (648 and 653 nm) and green (527 and 545 nm) bands are attributed to two-photon upconversion process, while that at the blue (450 and 475 nm) one is due to three-photon upconversion process [35,36]. In the case, the origin of the white-light emission can be well understood by a combination of two complementary colors (i.e., blue and yellow ones).

As predicted, the decreasing concentration of ${\mathrm{Yb}}^{3+}$ dopants predominately favors the green emission, whereas at higher concentrations, both the cross-relaxation and concentration quenching effects facilitate the red emission in the co-doped system [Fig. 6(c)] [16,19]. Thus, the yellow one happens as a result of the precise control of ${\mathrm{Yb}}^{3+}$ doping concentration. Clear evidence of yellow light [Figs. 6(c), 6(d), and 6(f)] appears from the excited ${\mathrm{CaF}}_2\text{:}{\mathrm{Yb}}_x{\mathrm{Er}}_1\text{-}{\mathrm{SiO}}_2$ ($x=35–45$) samples, of which the corresponding color coordinates locate at around (0.484, 0.510) in the Commission Internationale de l’Éclairage (CIE) 1931 chromaticity diagram. As the increase of the ${\mathrm{Yb}}^{3+}$ dopant, the CIE color points gradually change from (0.323,0.662) to (0.584,0.401) with the emitting color evolving from green through yellow to red. Note that the low green-to-red ratio is necessary to the generation of yellow light. In addition, the relative intensity of blue light was maximized in the ${\mathrm{CaF}}_2\text{:}{\mathrm{Yb}}_{35}{\mathrm{Tm}}_{1.5}\text{-}{\mathrm{SiO}}_2$ film sample. A further increase of ${\mathrm{Tm}}^{3+}$ [above 1.5 mol.%, Fig. 6(a)] and ${\mathrm{Yb}}^{3+}$ [above 35% mole fraction, Fig. 6(b)] concentration leads to the decrease of blue light, which is largely owing to the concentration quenching effect [20,22].

Then, the possibilities to support white-light emission from the luminescent films were studied from a series of ${\mathrm{CaF}}_2\text{:}{\mathrm{Yb}}_{35}{\mathrm{Tm}}_{1.5}{\mathrm{Er}}_y\text{-}{\mathrm{SiO}}_2$ ($\mathit{y}=0.1–2.0$) film samples [Figs. 6(e) and 6(f)]. Obviously, ${\mathrm{CaF}}_2\text{:}{\mathrm{Yb}}_{35}{\mathrm{Tm}}_{1.5}{\mathrm{Er}}_{0.5}\text{-}{\mathrm{SiO}}_2$ film generates the resulting white light, with the color coordinate at (0.329, 0.331), excited by a CW 980 nm laser at a power density of $20\mathrm{W}{\mathrm{cm}}^{-2}$, which is very close to that of the white point CIE standard white illuminant coordinates (0.33, 0.33). Furthermore, as the ${\mathrm{Er}}^{3+}$ concentration increases from 0 to 2.0% (mole fraction), Fig. 6(f) presents the sequence of the color evolution from blue through white to yellow. More importantly, a large range of the $\mathrm{CIE}(x,y)$ color space is enclosed by three elementary color samples, involving red (${\mathrm{CaF}}_2\text{:}{\mathrm{Yb}}_{95}{\mathrm{Er}}_1\text{-}{\mathrm{SiO}}_2$), green (${\mathrm{CaF}}_2\text{:}{\mathrm{Yb}}_{15}{\mathrm{Er}}_1\text{-}{\mathrm{SiO}}_2$), and blue (${\mathrm{CaF}}_2\text{:}{\mathrm{Yb}}_{35}{\mathrm{Tm}}_{1.5}\text{-}{\mathrm{SiO}}_2$) ones. This is comparable to that of the standard RGB (sRGB) benchmark, implying that this combined strategy by controlling the relative intensity of three elementary color bands should permit fluorescence of the ${\mathrm{Ln}}^{3+}$-doped system spanning the full-visible color range. Hence, this luminescent glass film would be an excellent lasing candidate for a CMOS-compatible device.

To confirm this assumption, the ${\mathrm{CaF}}_2\text{:}{\mathrm{Ln}}^{3+}$ NCs-doped ${\mathrm{SiO}}_2$ microtoroid cavity is fabricated upon a silicon substrate with a 2 μm thick oxide layer. A combination of a three-step patterning process, a spin-coating operation, and a subsequent laser-induced reflow process is employed to obtain the ${\mathrm{CaF}}_2\text{:}{\mathrm{Ln}}^{3+}$ NCs-doped ${\mathrm{SiO}}_2$ microtoroid cavity with ultrasmooth surface [17,37,38]. The fabrication details are depicted in Fig. 7(a). Initially, a thin negative photoresist layer (AZ2020) was spin-coated (4000 r/min, 60 s) onto the thoroughly cleaned oxidized silicon substrate and patterned using a mask aligner (Sichuan Nanguang HC 25). Through the standard photolithography (SVC Model H94-25C) and inductively coupled plasma (ICP, Oxford Plasmalab System100 ICP 180) etching processes, the circular pattern was transferred from initial mask, to negative photoresist, to silica layer, and finally to silicon substrate in sequence. In the ICP process, ${\mathrm{C}}_4{\mathrm{F}}_8$ and ${\mathrm{SF}}_6$ gases were specifically selected for the purpose of anisotropic silica etching and isotropic silicon etching, respectively. Consequently, a silica microdisk holding by the underlying silicon pillar was achieved. After removing the photoresist layer, the precursor containing ${\mathrm{Ln}}^{3+}$ ions was spin-coated onto the preformed ${\mathrm{SiO}}_2$ substrate, followed by an optimized four-stage calcination. Finally, an additional ${\mathrm{CO}}_2$ laser (CWQ800K)-assisted reflow was conducted to melt the microdisks along the periphery and therefore smooth the rims. The typical SEM images of such surface-tension-induced ${\mathrm{CaF}}_2\text{:}{\mathrm{Yb}}_{35}{\mathrm{Tm}}_{1.5}{\mathrm{Er}}_{0.5}\text{-}{\mathrm{SiO}}_2$ microtoroids are shown in Figs. 7(b)–7(d), where the microtoroid cavity well inherits the initial shape of pure silica microdisk (Fig. 8) after going through several pattern transferring processes. It is found that the overall disk diameter shrinks to form an 86 μm microtoroid. The height of the silicon pillar supporting the microtoroid was around 24 μm. No obvious variances in size and shape can be detected from those microtoroid cavities as long as the pumping power of the ${\mathrm{CO}}_2$ laser is fixed at a constant value.

The microcavities were first characterized in the 1550 nm band with a tapered fiber waveguide, as shown in the left inset of Fig. 9(a) [44,49]. A concise survey on the dependence of the $\mathit{Q}$ factor over a range of the coating thickness was conducted by recording the calculated $\mathit{Q}$ values. In this study, the thickness of the glass film could be determined by the spin-coating speed and repetition times, and a minimum layer thickness to achieve high $\mathit{Q}$ microtoroid resonator is found to be around 66 nm. From the right inset of Fig. 9(a), the loaded $\mathit{Q}$ of the ${\mathrm{CaF}}_2\text{:}{\mathrm{Yb}}_{35}{\mathrm{Tm}}_{1.5}{\mathrm{Er}}_{0.5}\text{-}{\mathrm{SiO}}_2$ microtoroid is as high as $5.04\times {10}^5$ in air through the equation $\mathit{Q}=\lambda /\delta \lambda$, where $\lambda$ is the resonant wavelength and $\delta \lambda$ is the full width at half-maximum (FWHM) of the transmission peak [34–36]. Obviously, the reflow operation facilitates the formation of an ultrasmooth cavity surface, which is consistent with the observations from the SEM images in Fig. 7.

The lasing characterization of the as-prepared microtoroid [Figs. 9(b)–9(e)] was performed by the free-space excitation. Generally, a pulsed 980 nm laser (6 ns, 10 Hz) from a 355 nm frequency-tripled $\mathit{Q}$-switched Nd:YAG laser (Continuum Surelite II-10) excited optical parameter oscillator (Continuum Horizon I) was directly focused on the top surface of the microtoroid. The emission from the edge of the resonator was collected by an optical fiber and recorded by an iHR-320 monochromator attached with a photomultiplier tube. To determine the onset of lasing action, Fig. 9(b) presents the RGB emission intensity over a range of power density. Clearly, the light-light curves exhibit S-like shapes, implying the transition from spontaneous emission through amplified spontaneous emission to lasing oscillation. Accordingly, the threshold values ($P_{\mathrm{th}}$), as reflected by the second kink in Fig. 9(b), are found to be as low as 1.99, 1.89, and $2.10\mathrm{mJ}{\mathrm{cm}}^{-2}$ for the red, green, and blue emissions, respectively. These similarly low threshold values verify the high $\mathit{Q}$ of such an active microtoroid cavity. As expected, the slopes (i.e., between the $P_a$ and $P_{\mathrm{th}}$) of the blue emission are larger than the red and green counterparts, while those of the red and green emissions share a similar value. This is reasonable since the blue emission and the green/red emissions are due to three-photon and two-photon upconversion process, respectively.

The spectral evolution of RGB emissions with the increase of pumping power is depicted in Figs. 9(c)–9(e). Only broadband emission peaks are observed at low power density. Then, periodic peaks maxima at 653, 545, and 475 nm sequentially occur and quickly dominate the emission spectra as the pumping power increases. Owing to the occurrence of stimulated emission, the FWHMs of these peaks collapse to less than 0.8 nm. We further investigate the mode spacing of the RGB bands. At a pumping level above $\sim 2\mathrm{mJ}{\mathrm{cm}}^{-2}$, the mode spacing, $\mathrm{\Delta }\lambda$, can be clearly observed from the high-resolution spectra, which are recorded at around 0.56, 0.75, and 1.0 nm for blue, green, and red emissions, respectively. The calculated $\mathrm{\Delta }\lambda$ can be taken from the equation $\mathrm{\Delta }\lambda ={\lambda }_0^2/n_{\mathrm{eff}}\pi d$, where ${\lambda }_0$, $n_{\mathrm{eff}}$, and $d$ are the emitting wavelength, the group refractive index, and the diameter of the microtoroid, respectively. Typically, for $\lambda =653\mathrm{nm}$, $n_{\mathrm{eff}}=1.45$ at 653 nm, and the diameter of the microtoroid cavity $d=86\mathrm{\mu m}$, the theoretical $\mathrm{\Delta }\lambda$ is around 1.08 nm, which is in reasonable agreement with the measured one. Furthermore, the inset of Fig. 9(e) presents the numerical simulation of the excited whispering gallery modes (WGMs) at the wavelength of 653 nm in 2D geometry. This is a clear indication of the formation of WGMs via total internal reflection along the interface of the volume, which is consistent with the measured results [46]. All these features unambiguously demonstrate that our microtoroid cavities can simultaneously support RGB lasing through the resonance of WG modes.

As has been previously realized, the efficiency of upconversion lasing, especially that at higher lying states, is suppressed due to the depletion of energy related to the high nonradiative recombination loss of ${\mathrm{Ln}}^{3+}$ ions in pure silica [23,24]. In sharp contrast with those ${\mathrm{SiO}}_2\text{:}{\mathrm{Ln}}^{3+}$ microcavity, our device fabrication allows ${\mathrm{Ln}}^{3+}$ ions to disperse in the ${\mathrm{CaF}}_2$ crystalline host lattice, which were selectively encapsulated in silica microtoroid as the gain medium. In this study, both the ultrasmall size and uniform distribution of NCs ensure a high $\mathit{Q}$ factor of the microtoroid ($>{10}^5$), implying strong light field confinement in the small volume. The doping density of ${\mathrm{Ln}}^{3+}$ ions can be increased to reach a high optical gain, whereas the clustering and detrimental interactions of ${\mathrm{Ln}}^{3+}$ ions are effectively suppressed due to the firmly confined ${\mathrm{Ln}}^{3+}$ in cubic ${\mathrm{CaF}}_2$ lattice and the uniform distribution of NCs in the matrix [32,50]. Multiphoton upconversion processes are also improved for the inherent low phonon energy of the ${\mathrm{CaF}}_2$ host [39,40,45]. Thus, an efficient white-light upconversion lasing was realized from the as-prepared ${\mathrm{CaF}}_2\text{:}{\mathrm{Ln}}^{3+}\text{-}{\mathrm{SiO}}_2$ microtoroid cavities.

Further evidence for stable white-light lasing behavior in the ${\mathrm{CaF}}_2\text{:}{\mathrm{Yb}}_{35}{\mathrm{Tm}}_{1.5}{\mathrm{Er}}_{0.5}\text{-}{\mathrm{SiO}}_2$ microtoroid can be seen in Fig. 10. Figure 10(a) presents the measured $P_{\mathrm{th}}$ and $1/\mathrm{\Delta }\lambda$ of the RGB emissions as a function of $d$. As anticipated, the $P_{\mathrm{th}}$ values decrease as the cavity size increases [34–36,51,52]. This can be defined with the equation $P_{\mathrm{th}}\sim \exp (-\alpha d)$, where $\alpha$ is the curvature loss [35]. In addition, the $\mathrm{\Delta }\lambda$ of the RGB bands remains inversely proportional to $d$, and the variations can be linearly fitted by the blue lines in Fig. 10(a) [34–36]. Figure 10(b) exhibits the simultaneous RGM emission spectra (i.e., above $P_{\mathrm{th}}$) of each microtoroid with different $d$ values (i.e., 56, 86, and 117 μm, respectively). The corresponding color points, involving (0.324, 0.354), (0.327, 0.355), and (0.377, 0.324), are depicted in Fig. 10(c). Furthermore, the CIE 1931 chromaticity diagram in Fig. 10(d) shows the calculated color coordinates, which locate at (0.327, 0.355), (0.335, 0.365), and (0.346, 0.361), respectively, of the excited microtoroid ($d\approx 86\mathrm{\mu m}$) at different power densities. Intriguingly, all these points fall within the white-light region with a very tiny shift when the pumping density is above $2.25\mathrm{mJ}{\mathrm{cm}}^{-2}$.

This white-light microlaser is intrinsically different from the previous reports [9,12,13]. In conventional designs, the thresholds of RGB microlasers are quite different. As a result, the ratios between different colors change dramatically, strongly affecting the stability of the white-light laser. In contrast, the color coordinates of our white-light lasing roughly remain stable (i.e., above $P_{\mathrm{th}}$) even with the increase of pumping power. On the one hand, such stability is reasonable owing to the conspicuous difference between unsaturated upconversion [13–15] dominating in photoluminescence and gain saturation [17,22,51,52] dominating in lasing emission. Generally, the double-logarithmic dependence of blue emission (i.e., three-photon upconversion process) and red, green emissions (i.e., two-photon upconversion process) on the pumping power can be well fitted by a linear equation, $\lg I_{\mathrm{em}}=k\lg I_{\mathrm{ex}}$, where $I_{\mathrm{em}}$ is the integrated output intensity, $I_{\mathrm{ex}}$ is the pumping power, and $k$ is the slope representing the absorbed excitation photons [13,14]. For PL, the $k$ values of the RGB emissions are very different due to distinct multiphoton upconversion processes. Thus, the resulting luminescent color related to the proportions of RGB emissions changes over a wide range with the increasing irradiation. On the contrary, the RGB emissions dominating at higher power intensity (i.e., above $P_{\mathrm{th}}$) share slopes of around 1 as the onset of lasing emission [17,22,51,52]. On the other hand, the simultaneous RGB lasers exhibit similar thresholds due to the highly confined WG modes. Consequently, the ratio of R:G:B emissions remains roughly the same for a wide power range other than negligible changes in the proportions of the blue and red color bands in the total integrated intensity. All these well explain the stability of our high-$\mathit{Q}$ white-light laser on a chip.

In conclusion, we have demonstrated a high-$\mathit{Q}$ (i.e., exceeding ${10}^5$) ${\mathrm{CaF}}_2\text{:}{\mathrm{Ln}}^{3+}$ UCNCs-doped ${\mathrm{SiO}}_2$ toroidal microcavity on a silicon chip. The key to our successful demonstration is a controlled growth approach that permits a uniform distribution of the ultrasmall ${\mathrm{CaF}}_2\text{:}{\mathrm{Ln}}^{3+}$ NCs in ${\mathrm{SiO}}_2$ matrix, and therefore the development of a high-$\mathit{Q}$ active microtoroid. By stoichiometric doping of ${\mathrm{Ln}}^{3+}$ ions, such microlasers are capable of generating simultaneous RGB lasing with similar thresholds and robust white-light upconversion lasing from a single gain medium. Meanwhile, as the onset of highly confined WGM lasing, the increasing rates of the RGB emissions above $P_{\mathrm{th}}$ are roughly the same, thus resulting in the similar proportions of the RGB light in the total integrated intensity, which determines the excellent stability of the resultant on-chip white-light lasing. Benefiting from the doping flexibility, such high-$\mathit{Q}$ chip-integrated lasers that could lase at any color of interest with high stability see critical applications in many fields.

Reagents. All chemical regents, involving ytterbium (III) acetate hydrate (99.9%), thulium (III) acetate hydrate (99.9%), erbium (III) acetate hydrate (99.9%), calcium acetate hydrate (99.0%), trifluoroacetic acid (TFA, 99.0%), TEOS (99.99%), ethanol absolute, acetic acid (AA, 99.7%), were used as received without further purification.

Synthesis of$\mathbf{C}\mathbf{a}{\mathbf{F}}_{\mathbf{2}}\text{:}{\mathbf{Ln}}^{\mathbf{3}+}$nanocrystals-doped silica film. A series of ${\mathrm{CaF}}_2\text{:}{\mathrm{Yb}}^{3+}/{\mathrm{Tm}}^{3+}/{\mathrm{Er}}^{3+}$ ($x$, $y$, $z$ represent mole fractions of individual ions) nanocrystals-doped silica nanofilm was synthesized by solgel and spin-coating procedure [37,38]. In a typical synthesis, raw TEOS was treated in dilute ethanol ($\mathrm{TEOS}\text{:}\mathrm{ethanol}=1\text{:}1$, volume ratio) with required amounts of AA as catalytic agent under vigorous stirring at 70°C. TEOS, de-ionized ${\mathrm{H}}_2\mathrm{O}$, and AA were added in the molar ratio of 1:10:0.5. Stoichiometric amounts of the lanthanide (III) and calcium acetate salts were dissolved in TFA. The molar ratio of metal ions to ${\mathrm{F}}^{-}$ was 1:10, and that of metal ions to silica was 8:92. The mixture was heated at 70°C and magnetically stirred for 90 min for a clear solution. After cooling down to room temperature, the solution was added to the silica precursor dropwise. The resultant solution was stirred for 1 h to form a sol and then spin-coated on an oxidized silicon substrate. Finally, the coatings were annealed by the following steps: (i) dried from room temperature to 110°C at a heating rate of $0.25°\mathrm{C}{\min }^{-1}$, and staying for 20 min; (ii) heated to 260°C at a rate of $0.5°\mathrm{C}{\min }^{-1}$; (iii) heated to 400°C at a rate of $1°\mathrm{C}{\min }^{-1}$; and (iv) heated to 800°C at a rate of $5°\mathrm{C}{\min }^{-1}$.

Characterization. The formation of ${\mathrm{CaF}}_2\text{:}{\mathrm{Ln}}^{3+}\text{-}{\mathrm{SiO}}_2$ nanofilm was confirmed by scanning electron microscopy (Hitachi Model S-4700, 10–30 kV), transmission electron microscopy (FEI/Philips Tecnai 12 BioTWIN, 120 kV), high-resolution TEM (JEOL-JEM 2100F, 200 kV), and X-ray diffraction (Bruker AXS D2). The thickness of the silica film was analyzed by a thickness gauge (Dektak 150). The thermogravimetry/differential thermal analysis was carried out on a Netzsch STA 449 F5 thermal analyzer. The ultraviolet-visible absorption and transmission spectra of the film were performed on a Shimadzu UV-2550 spectrometer. High-resolution emission spectra were recorded by an iHR-320 monochromator (Horiba) attached with a photomultiplier. The resolution of the spectrometer is 0.06 nm.

Acknowledgment. The authors also acknowledge support from the Shenzhen Engineering Laboratory on Organic-Inorganic Perovskite Devices.