Since the 1980s, spintronics has been intensely researched for the next generation of data operation and storage in information technology, because spintronics could offer an exciting prospect of combining the semiconductor-based information operation (logic) and magnetic based data storage (memory) into the same device1,2,3,4. However, the practical realization of spintronic devices remains challenging due to the lack of functional materials displaying both, semiconducting and high-temperature magnetic ordering properties5,6,7. Historically, the research of ferromagnetic semiconductor materials has been primarily focused on the development and characterization of inorganic dilute magnetic semiconductors (DMSs), such as p-type Mn-doped II–VI8, III–V9, and IV–VI semiconductors10,11. Recently, intrinsic ferromagnetism has been demonstrated in layered inorganic two-dimensional (2D) materials, such as chromium germanium telluride (Cr2Ge2Te6)12 and chromium triiodide (CrI3)13. Despite enabling long-range electronic coupling via short metal-metal bonds in these inorganic magnets14, the low magnetic ordering temperature (e.g., Tc = 159 K for GaMnAs, 61 K for Cr2Ge2Te6, and 68 K for CrI3, see Supplementary Table 1), together with the limited chemical tunability present major drawbacks for their practical applications.

Metal-organic frameworks (MOFs) based on crystalline coordination polymers have been regarded as promising organic magnets15. These materials hold promise because the magnetic behavior can be tailored at the molecular scale through tuning their constituent organic ligands, metal centers (spin carriers) and even filling functional guest molecules in porous MOFs16. Magnetic ordering at low temperature (Tc < 180 K, seen in Supplementary Table 2) has been already proven in bulk (3D) MOFs14. However, most of them behave as electrical insulators (σ < 10−12 s m−1) because of the localized molecular orbitals and minimal band dispersion in their intrinsic electronic structures, thus hampering their potential applications in spintronics/electronics. Taking advantage of the porous structure, hosting guest molecules (like tetracyanoquinododimethane) in MOFs has recently been shown to significantly enhance the electrical conductivity17, due to the charge transfer of guests to the metal nodes or organic linkers. In addition, linking redox-active ligands, such as 2,5-dihydroxybenzoquinone14,18, pyrazine19, could also generate long-range charge transport and strong magnetic exchange, leading to conductivities as high as 20 s m−1 and magnetic ordering. On the other hand, since the first report in 201220, conjugated MOFs with layered structures have been developed displaying high electrical conductivities (up to 105 s m−1). This class of materials has typically been constructed by linking N, O, or S ortho-disubstituted benzene, triphenylene or coronene ligands with transition metal ions, and featured with full π-d conjugation in 2D planes20,21,22,23,24,25,26,27. Furthermore, density functional calculation (DFT) has predicted that strong planar π-d hybridization combined with large magnetic anisotropy could generate magnetic ordering at room temperature28. However, the simultaneous realization of room temperature spontaneous magnetization and semiconducting behavior in a MOF has not been experimentally demonstrated to date.

Here, we report a conjugated K3Fe2[PcFe–O8] MOF with square lattice geometry and interplane van der Waals (vdW) layer-stacking structure by employing (2,3,9,10,16,17,23,24-octahydroxy phthalocyaninato) iron (PcFe-OH8) as ligand and square planar iron-bis(dihydroxy) complex as linkage as well as K+ as the counter ions. The resultant K3Fe2[PcFe–O8] exhibits a typical semiconducting behavior with a room temperature carrier mobility of 15 ± 2 cm2 V−1 s−1 as determined by time-resolved Terahertz spectroscopy (TRTS). Moreover, this K3Fe2[PcFe–O8] presents magnetic correlation in nanoscale crystallites, which overall shows superparamagnetism with a broad distribution of blocking temperatures. DFT calculations reveal that the strong hybridization of the π-d orbitals and the high concentration of Fe spins (0.7 μB per Fe site) in K3Fe2[PcFe–O8] contribute to its high-temperature magnetic exchange interactions. By further optimizing the crystalline quality and increasing the crystallite size, our work presents the possibility to achieve room temperature ferromagnetism in a semiconducting layered MOF, highlighting the potential for developing a new generation of MOFs-based spintronics.


Synthesis and structural analysis of K3Fe2[PcFe–O8]

The K3Fe2[PcFe–O8] MOF with potassium cation (K+) as counter ions was synthesized through the coordination reaction between PcFe-OH8 and iron (II) acetate (Fe(OAc)2) in the presence of potassium acetate (KOAc) (Fig. 1a, Supplementary Fig. 1, see Supplementary Methods). Fourier transform infrared spectra (FTIR, Supplementary Fig. 2) showed that O–H stretching bonds (~3300 cm−1) disappeared in K3Fe2[PcFe–O8], indicating the successful coordination reaction between the –OH groups and Fe ions. Besides, new peaks centered at 534 and 468 cm−1 can be assigned to O–Fe–O stretching bonds. Energy-dispersive X-ray (EDX, Supplementary Figs. 34) spectroscopy and X-ray photoelectron spectroscopy (XPS, Supplementary Fig. 5) were carried out to analyze the composition of K3Fe2[PcFe–O8], which revealed the presence of C, N, O, Fe and K in the MOF sample. The core level spectrum of Fe (2p) contains peaks at 710.8, 713.4, 724.2, 726.6 eV, attributable to the Fe(II) 2p3/2, Fe(III) 2p3/2, Fe(II) 2p1/2, and Fe(III) 2p1/2, respectively, revealing that two kinds of iron ions (Fe2+/Fe3+) populate the K3Fe2[PcFe–O8] samples (Supplementary Fig. 5). The corresponding Fe2+/Fe3+ ratio was inferred to be ~2/1. The O (1s) spectra exhibit a Fe–O–C peak at a binding energy of 531.3 eV, further demonstrating the coordination reaction between O and Fe atoms. The signal referring to K in K3Fe2[PcFe–O8] XPS spectrum originates from the counter ions balancing the—otherwise—negatively charged system. Element analysis and TGA measurements (Supplementary Fig. 6) further define the chemical formula of the layered MOF as K3Fe2[PcFe–O8]·2.2H2O (C32H12.4Fe3K3N8O10.2; named as K3Fe2[PcFe–O8] for short). In order to exclude the effect of water adsorption on the surface area and porosity, K3Fe2[PcFe–O8] was treated with supercritical carbon dioxide (CO2) drying29, after which low-pressure N2 sorption was measured at 77 K. The obtained spectra (Supplementary Fig. 7) exhibit a superposition of type II and IVa isotherms30 which is typical for nanosized porous crystals. The Brunauer-Emmett-Teller surface area of K3Fe2[PcFe–O8] is estimated as 206 m2 g–1, with the counter ions K+ incorporated in the pores.

Fig. 1
figure 1

Synthesis and characterization of layered K3Fe2[PcFe–O8] MOF. a Schematic illustration for the synthesis of Fe2[PcFe–O8] framework with iron ions and organic PcFe-OH8 linkers connected by coordination bonds (light cyan: C; blue: N; light pink: O; orange: Fe3+ in the phthalocyanine ring; green: Fe2+ in the linkage; H atoms and K+ counter-ions omitted for clarity). The interval between layers is about 3.3 Å along the c-axis; b Experimental PXRD pattern (organge curve) and that of calculated (DFT) AA-serrated stacking structures (dark cyan curve) of K3Fe2[PcFe–O8]; c HRTEM image of K3Fe2[PcFe–O8]. Scale bar: 5 nm. Inset: corresponding FFT analysis; d Normalized Fe K-edge XANES spectra of K3Fe2[PcFe–O8], Fe foil, FeO, Fe2O3, and PcFe-OH8. Insert: enlarged pre-edge region in Fe K-edge XANES spectra; e, Fourier transformation EXAFS spectra at Fe K-edge of Fe2[PcFe–O8] with Fe2O3 and PcFe-OH8 as contrast

Powder X-ray diffraction (PXRD) measurements with Cu Kα irradiation (λ = 1.54 Å) for K3Fe2[PcFe–O8] revealed a highly crystalline structure with prominent (100), (200), (300) and (001) peaks at 2θ = 4.9°, 9.8°, 14.7°, 26.6°, respectively (Fig. 1b and Supplementary Fig. 8). We combined DFT calculation to optimize the atomic structure of single layer MOF and then simulate different stacking arrangements of layer-stacked K3Fe2[PcFe–O8], e.g., AA, AA-inclined, AA-serrated, and AB stacking modes (Fig. 1b and Supplementary Fig. 9)31. The DFT results show that AA-serrated stacking geometry is energetically favored and its calculated PXRD pattern agrees well with the experimental result. The corresponding unit cell parameters with the P1 group are a = b = 18.1 Å, α = 89.5°, β = 89.8°, and γ = 89.3°, indicating long-range ordering along the ab plane. The weak and broad peak at 26.6° corresponding to the (001) facet suggests a 3.3 Å interval between layers in c-axis normal to the 2D plane. Additionally, the average size of crystalline domains in K3Fe2[PcFe–O8] was calculated to be ~36 nm according to the Scherrer’s equation32. Transmission electron microscopy (TEM) results reveal rectangle-shaped crystals with sizes in the range of 10–100 nm (Supplementary Figs. 10 and 11). High-resolution TEM image (Fig. 1c) and the corresponding fast Fourier transform (FFT) image (inset Fig. 1c) manifest a K3Fe2[PcFe–O8] sample architecture based on a square lattice with a pore size of ~1.75 nm.

Synchrotron powder X-ray adsorption spectroscopy (XAS) was employed to further analysis the chemical state and coordination geometry in K3Fe2[PcFe–O8]. Reference samples, e.g., Fe foil, FeO, Fe2O3 as well as the precursor PcFe-OH8 were also investigated by XAS. Fig. 1d shows the X-ray adsorption near-edge structure (XANES) spectra at K edge of all the samples. The Fe K-edge of XANES in K3Fe2[PcFe–O8] exhibits a near-edge spectra similar to that of PcFe-OH8 monomers, but completely different from those of FeO, Fe2O3, and Fe foil. Besides, the pre-edge feature (magnified in Fig. 1d) in K3Fe2[PcFe–O8] originated from the transition of 1s core electrons to hybridized orbitals of Fe (3d) and ligands (p)33. The intensity of the pre-edge peak is more intense on the site symmetry where the iron atom is located34. Figure 1e displays the Fourier transform of the κ-weighted extended X-ray absorption fine structure (EXAFS) of K3Fe2[PcFe–O8] as well as the contrast samples. The EXAFS presents a predominant peak in K3Fe2[PcFe–O8], which is originated from the nearest-neighboring nitrogen or oxygen coordination shell around the Fe atoms35. Based on this peak, Fe–N(O) distance was calculated to be ~1.57 Å. From the shape and amplitude of the main peak in the magnitude of the FT spectra, it is obvious that the bonding environment in K3Fe2[PcFe–O8] is very close to that of the square planar geometry of PcFe–OH8, further suggesting Fe atoms connected with four N/O atoms in the layered MOF. However, due to the limitation of the XAS resolution, we cannot distinguish the differences between Fe–O4 and Fe–N4 coordination geometries in the MOF. Nevertheless, another contrast sample Fe2O3 clearly exhibits two different predominant peaks at ~1.44 Å and ~2.57 Å, which arise from Fe–O and Fe–Fe bonds, respectively (Fig. 1e). Therefore, the XANES and EXAFS spectra of K3Fe2[PcFe–O8] and the contrast experiments provide strong proof on the formation of square planar complexes via the coordination of PcFe–OH8 and Fe ions. Moreover, no metal oxides such as FeO and Fe2O3 were detected in the K3Fe2[PcFe–O8].

Compared with that of the constituent monomer PcFe–OH8, the UV–Vis spectrum of K3Fe2[PcFe–O8] powders in N, N-Dimethylformamide (DMF) reveal a red-shift of the Soret-band to ~390 nm (Supplementary Fig. 12), testifying to the extended π-conjugation36. Cyclic voltammetry (CV) curves of K3Fe2[PcFe–O8] indicate a narrow energy gap of about 0.63 eV for the MOF powders dispersed in acetonitrile (Supplementary Fig. 13, Supplementary Table 3)37.

Electronic structure of K3Fe2[PcFe–O8]

To estimate the electronic structures of K3Fe2[PcFe–O8], we performed the DFT calculations. The band structure of a monolayer K3Fe2[PcFe–O8] with optimized geometry shows relatively dispersed bands on both sides of the Fermi level, suggesting a semiconducting behavior in this configuration with a rather narrow bandgap of ~0.12 eV (Fig. 2a, enlarged figure seen in Supplementary Fig. 14). The computational approach used in this study (DFT + U and PBE exchange-correlation functional, seen in Supplementary Methods), usually underestimates the bandgap in semiconductors38. Besides, the bands near enough to the Fermi level could be easily thermally-populated with holes, indicating a typical p-type semiconducting behavior. The PDOS near the Fermi level exhibits considerable hybridization of the orbitals from Fe(d), C(p), O(p) and N(p), which demonstrates a high degree of π-d conjugation in the monolayer plane. The spin density iso-surface of K3Fe2[PcFe–O8] (Fig. 2b and Supplementary Fig. 15) reveals that the spin density is primarily localized on the Fe atoms in phthalocyanine cores, while the square planar Fe–O4 moieties are slightly polarized by the polarization of the delocalized π orbitals among Fe(d), C(p), O(p) and N(p)28. Notably, the oxidation state of Fe ions in the phthalocyanine ring is always higher than that in the Fe–O4 moieties with their relatively higher Fe 2p energy levels of 1.5–2.5 eV (Supplementary Fig. 16), which supports the coexistence of Fe(II)/Fe(III) determined by XPS.

Fig. 2
figure 2

Modeling of the electronic structures of K3Fe2[PcFe–O8]. a Band structure of a monolayer K3Fe2[PcFe–O8] with GGA + U correction shown on the left panel, dashed lines indicate the bands associated to spin up while the solid lines indicate the bands associated to spin down (the effective Coulomb (U) and exchange (J) terms reported in the Supplementary Information). The corresponding projected density of states (PDOS) for spin up and spin down are plotted on the right panel for Fe(d), C(p), O(p), and N (p) states; b Spin density iso-surface (pink solid iso-surface), at absolute spin-density |ρ ↑–ρ↓| = 0.05 electrons per Å3 of a monolayer K3Fe2[PcFe–O8], indicating that the spin density is mainly localized on the Fe ions (light cyan: C; blue: N; red: O; light red: Fe; gray: K+); c Calculated electronic band structure of multi-layered K3Fe2[PcFe–O8] with AA-serrated stacking mode; d Unit cell of two-layered K3Fe2[PcFe–O8] in AA-serrated stacking mode with ferromagnetic arrangement

We further investigated the electronic band structure and PDOS of multi-layered K3Fe2[PcFe–O8]. Fig. 2c displays a rather narrow band gap of ~0.10 eV in the AA-serrated stacking structure. In this case, Fig. 2d shows spin density iso-surfaces of two-layered K3Fe2[PcFe–O8] with unit-cell magnetization, in which the magnetic moments are also predominantly localized on the Fe atoms in phthalocyanine cores. Moreover, the magnetic ground state calculation of the AA-serrated stacking system reveals an exchange energy of Eex = 300 meV (Eex = EAFMEFM), which implies the possible intrinsic ferromagnetic ordering of K3Fe2[PcFe–O8] (Supplementary Fig. 17).

Charge transport and magnetotransport in K3Fe2[PcFe–O8]

The electrical conductivity of the K3Fe2[PcFe–O8] samples through van der Pauw method was derived to be 2 × 103 S m−1 at 350 K (measured in pellets with a thickness of ~0.59 mm, Supplementary Fig. 18).25 Variable-temperature conductivity measurement presented a non-linear increase of electrical conductivity from 140 to 350 K, indicating a typical semiconducting behavior of the K3Fe2[PcFe–O8] sample (Fig. 3a and Supplementary Fig. 19). The electrical conductivity plotted versus reciprocal temperature (T−1) indicates two thermally activated contributions to the conductivity of K3Fe2[PcFe–O8]39. From the fitting results (inset in Fig. 3a), the activation energy in a high activation region was calculated to be ~261 meV, while that in a low-temperature activation region was ~115 meV (Supplementary Fig. 20). We ascribe this thermally-activated hopping progress to the grain boundaries dominating the temperature dependence of conductivity in the bulk polycrystalline pellets, thus giving rise to the relatively large activation energy23,40. In addition, it needs to mention that redox state of the iron centers also contributes to the conductivity of the K3Fe2[PcFe–O8]. As contrast, after the sample was oxidized in the air for one week, the conductivity was 2 orders of magnitude lower than that before the oxidation at room temperature (Supplementary Fig. 21a). The Fe (2p) core level XPS spectrum evidences the partial oxidation of Fe2+ into Fe3+ with a ratio of ~1/3 in the oxidized K3Fe2[PcFe–O8] (Supplementary Fig. 21b).

Fig. 3
figure 3

Charge transport and magnetoresistance of K3Fe2[PcFe–O8]. a Variable-temperature electrical conductivity of compressed K3Fe2[PcFe–O8] pellets via van der Pauw method; Insert: Plot of electrical conductivity of K3Fe2[PcFe–O8] as a function of inverse temperature (T−1); b Room-temperature photoconductivity of K3Fe2[PcFe–O8] measured by optical pump-THz probe spectroscopy (800 nm excitation wavelength 300 μJ cm−2 and ~1 THz bandwidth probe; under nitrogen environment). Red line represents a bi-exponential decay function; c Real (black dots) and imaginary (red dots) components of the frequency-resolved complex conductivity (measured at 0.5 ps after photoexcitation, DC conductivity 0.14 S m−1); solid lines represent a Drude-Smith description of the data; d, Magnetic field dependence of the magnetoresistance by measuring the changes of the electrical resistance in an applied field (−6–6 T) at different temperatures

The Hall resistance (RHall) was further measured under magnetic field (H) based on the van der Pauw pattern (Supplementary Fig. 22), which displays a linear increase of RHall with H. The fitting slope of the Hall effect shows a p-type semiconducting behavior in K3Fe2[PcFe–O8] with the hole concentration of about 2.4 × 1014 cm−3. The corresponding Hall mobility was calculated as ~0.1 cm2 V−1 s−1. However, the carrier mobility from the Hall measurement is much lower than its intrinsic value, due to the contact resistance and the grain boundaries/gaps of K3Fe2[PcFe–O8] particles in the pressed pellets (Supplementary Fig. 23).

Next, contact-free time-resolved THz spectroscopy (TRTS) was employed to further address the photoconductivity and mobility nature in the K3Fe2[PcFe–O8] samples. Figure 3b shows the real part of the photoconductivity measured as a function of pump-probe delay for a ~210-micron thick sample (300 μJ cm−2, 800 nm pump excitation, ~1 THz probe center frequency and ~1 THz bandwidth). These dynamics indicate the ultrafast (sub-ps) formation of quasi-free carriers that undergo rapid trapping and/or localization within the samples. Figure 3c shows the real and imaginary conductivity components as a function of frequency at the peak conductivity (0.5 ps after pump excitation); the data can be well described by the Drude-Smith model (shown as solid lines in Fig. 3c, seen in Supplementary Methods)41. This model for carrier transport is commonly employed to describe free charge carrier motion constrained by backscattering (as expected in our polycrystalline samples where long-range DC conductivity is suppressed by grain boundaries). The best fit to the data by the DS model provides a scattering time of 53 ± 8 fs and a c parameter of −0.69 ± 0.02 (where c ranges between 0 and −1; with 0 and −1 denoting completely free and completely localized free charges, respectively). From these estimates, and considering an e-h averaged effective mass of m*~1.9m0 (obtained from DFT calculations), we obtain a high mobility estimate for the photo-generated free carriers at the peak conductivity of 15 ± 2 cm2 V−1 s−1 (a mobility comparison with thus far reported stacked coordination polymers seen in Supplementary Table 4). Magnetotransport as a function of temperature was measured in the samples. Figure 3d reveals the change in resistance with varying applied magnetic field at 150, 200, and 300 K. As evident from the data, the resistance decreases with increasing magnetic field due to the reduced spin scattering10, indicating the interaction between the magnetic moments and the carriers.

Magnetic properties of the K3Fe2[PcFe–O8]

To probe the magnetic behavior of the K3Fe2[PcFe–O8], field- and temperature-dependent magnetization measurements were performed using superconducting quantum interference device (SQUID) magnetometry. Figure 4a shows the variation of the magnetization with the applied magnetic field (M-H curves) measured at different temperatures from 5 to 350 K (enlarged figure shown in Supplementary Fig. 24). Notably, at 350 K, K3Fe2[PcFe–O8] retains a magnetic hysteresis with a saturation magnetization of ~0.13 μB per Fe-site. Upon cooling to 5 K, an obvious increase in the coercivity and the magnetization was observed. Particularly, the hysteresis still did not reach the saturation at 7 T with magnetization of ~0.7 μB per Fe-site at 5 K. These observations suggest the superparamagnetic nature in the sample. Nevertheless, the magnetic hysteresis loops at 300 K and 350 K point to ferromagnetic coupling within single K3Fe2[PcFe–O8] crystallites with Curie temperature higher than 350 K. This is consistent with the DFT calculations which reveal an exchange energy of around 300 meV and also with the negative magnetoresistance at 300 K (Fig. 3d).

Fig. 4
figure 4

Magnetic properties of the K3Fe2[PcFe–O8]. a Magnetic hysteresis loops obtained at different temperatures for K3Fe2[PcFe–O8]; b Zero-field-cooled (ZFC) and field-cooled (FC) magnetization for K3Fe2[PcFe–O8] in an applied DC magnetic field of 100 Oe; c, d 57Fe Mössbauer spectra of K3Fe2[PcFe–O8] measured at a temperature of 300 K and 25 K, respectively

The blocked superparamagnetic nature is also confirmed by the zero field cooled (ZFC) and field-cooled (FC) measurements (Fig. 4b). The FC magnetization in an external field of 100 Oe reveals a monotonic decrease upon increasing temperature. The corresponding ZFC magnetization at the same external field increases with temperature with a broad summit peak around 200 K, approaching the FC curve at around 300 K. The irreversibility between the ZFC and FC curves is due to blocked superparamagnetic clusters which have a broad distribution of blocking temperatures. This is the result of structural inhomogeneities in K3Fe2[PcFe–O8] (shown in the TEM images, Supplementary Figs. 10 and 11). The nanoscale K3Fe2[PcFe–O8] crystallites behave as single-domain superparamagnets. A splitting feature between the ZFC and FC curves is observed up to 300 K, which is also indicative of the persistence of magnetic couplings between the Fe spins in individual crystallites at or above 300 K. Notably, the remnant magnetization (Supplementary Fig. 25) does not decrease to zero at 350 K, further suggesting the presence of superparamagnetic particles with blocking temperatures even above 350 K. These results suggest that the short Fe–Fe distance in the c direction together with the strong coupling between Fe spins mediated by the fully delocalized π-electrons along the ab plane contributes to the persistence of spontaneous magnetization above room temperature. Compared with the Curie temperatures of thus far reported magnetic MOFs14,18,42,43,44, our work discovers the layered conjugated MOF exhibiting ferromagnetic coupling above room temperature (Supplementary Table 2). However, the polycrystalline nature of the samples (10–100 nm in domain size) renders the observed superparamagnetism, which could be addressed by improving the sample preparation, such as to obtain single crystals in the future.


To further understand the magnetic mechanism in K3Fe2[PcFe–O8], we investigated the local environment of Fe atoms by 56Fe Mössbauer analysis. At 300 K, two kinds of Fe sites are resolved in the sample as shown in Fig. 4c. The narrow quadrupole doublet with an isomer shift of 0.378(4) mm s−1 and an electric quadrupole splitting ΔEQ = 0.843(6) mm s−1 indicate the Fe(III)-high spin oxidation state and provides the magnetic moments (Supplementary Table 5). The broad subspectrum assigned to Fe site 2 extends over the full velocity range with an isomer shift of 0.242(10) mm s−1 and a slowly fluctuating magnetic hyperfine field (Blume model). The fit results in a magnetic hyperfine field of 30(5) T with a fluctuation frequency of 150(50) MHz. Such a subspectrum is often observed in nanoscale superparamagnetic particles above the blocking temperature. In contrast, the 57Fe Mössbauer spectrum demonstrates the only presence of one Fe(III) state in the monomer PcFe–OH8 (Supplementary Fig. 26), which reveals that phthalocyanine ligands contribute to partial Fe site 1 while the Fe site 2 originates from the linkages in K3Fe2[PcFe–O8] MOFs. This observation is consistent with the ferromagnetic exchange within the individual clusters45.

Upon cooling from room temperature to 25 K, three Fe sites are observed in K3Fe2[PcFe–O8] (Fig. 4d) as expected. The spectrum at 25 K exhibit a sharp quadrupole splitting (Fe site A; isomer shift 0.35(1) mm s−1; ΔEQ = 0.94(1) mm s−1), which is in consistent with a Fe(III) high spin states (Fig. 4d), while the Fe site C is in consistent with that observed in the monomer (Supplementary Fig. 26). Besides, Fe site B shows a magnetic hyperfine splitting, the sixtet of which is consistent with the subspectrum observed in ferromagnetic nanoparticles. The Mössbauer measurements suggest the superparamagnetic nature of the sample. In the superparamagnetic state, the magnetization direction of nanoparticle fluctuates among the easy axes of magnetization when there is no external magnetic field. The relaxation time depends on the size of the particles and the temperature. Therefore, magnetically, three sites are observed at 25 K because some particles are large enough to have a longer relaxation time, which contributes to the sixtet, while the remained randomly oriented particles are superparamagnetic (red site A).

Therefore, based on the 56Fe Mössbauer analysis as well as the support from the XPS (Supplementary Fig. 5c) and spin density distribution calculation (Fig. 2d), we can infer that the superparamagnetism in our K3Fe2[PcFe–O8] readily origins from the polycrystalline feature by regarding to the varied crystalline domain sizes. In addition, it is proposed that the magnetic coupling is induced by indirect exchange interaction between the localized iron spins via the highly delocalized π electrons along the fully conjugated backbones10,24, thus enabling strong hybridization between the d/p orbitals of Fe, the phthalocyanine core, and the Fe–O4 nodes (seen in Fig. 2c). As mentioned before, the calculated magnetic ground states of AA-serrated K3Fe2[PcFe–O8] present an energy-favorable ferromagnetic coupling with a positive exchange coupling energy of 300 meV (Supplementary Fig. 17), which further validates this conclusion. Nevertheless, it should be noted that further experiments are required to improve the crystalline quality to obtain ferromagnetic MOFs at or even above room temperature and sophisticated characterizations should be done to unambiguously exclude the possibility of magnetic secondary phases. Our work is expected to encourage more physical researches on magnetic and semiconducting properties of layered conjugated MOFs.

In summary, we have demonstrated a phthalocyanine-based MOF (K3Fe2[PcFe–O8]) with layered structures, which exhibits a p-type semiconducting behavior with a high mobility of ~15 cm2 V−1 s−1 at 300 K and blocked superparamagnetic nature up to 350 K (the highest temperature we have measured). DFT calculations indicate that the ferromagnetic ground state benefits from the strong hybridization between the d-p orbitals of iron, the phthalocyanine core, and the iron-bis(dihydroxy) nodes. Our work highlights layered conjugated MOFs as a class of semiconducting materials for potential spintronics applications.



All the starting materials were purchased from commercial suppliers, such as Sigma-Aldrich, TCI and abcr GmbH. Unless otherwise stated, all the chemicals were used directly without further purification. The reactions were performed with the standard vacuum-line and Schleck techniques under nitrogen. Colum chromatography was performed using the silica gel.

(2,3,9,10,16,17,23,24-octahydroxy phthalocyaninato)Fe (PcFe–OH8) was synthesized according to the previous reported literature with minor modification (see Supplementary Information)46.

All the reactions were performed under vacuum using the Schlenk line technique.

Typical synthesis of K3Fe2[PcFe–O8]

In a 25 mL tube, PcFe-OH8 (20 mg, 0.029 mmol), iron (II) acetate (9.4 mg, 0.054 mmol) were mixed together with KOAc (26 mg, 0.26 mmol) in N-methyl-2-pyrrolidone (NMP)/H2O (3 mL/1 mL, v/v). The mixture was sonicated for 1 min, after which the system was degassed with freeze-pump-thaw methods for three times, sealed under high vacuum and kept at 423 K (150 °C) for 3 days. Then, the mixture was cooled to room temperature, filtrated and soaked in degassed H2O for 1 day. After washed by acetone and filtration, the layered K3Fe2[PcFe–O8] MOF was dried under vacuum at 40 °C as dark black powder. Yield: 88.4%. Element analysis and TGA measurements (Supplementary Fig. 6) define the chemical formula of the MOF as K3Fe2[PcFe–O8]·2.2H2O (C32H12.4Fe3K3N8O10.2; found Fe, 17.3; C, 40.16; H, 1.38; N, 11.83; K, 12.13; O, 17.05; named as K3Fe2[PcFe–O8] for short).

Variable-temperature conductivity measurements

The pressed pellets were prepared by adding 25 mg samples (heated at 100 ℃ under vacuum overnight) onto a polymer film in an 8 mm inner diameter split sleeve pressing under the applied pressure of 1 GPa at 100 °C. After the temperature cooled down to room temperature, the pressed pellet was taken out and the thickness was measured. Then, four probes of silver wires were placed onto the top of the pressed pellets using conductive silver plastic. The probe was transferred onto the probe station with thermally conductive and electrically insulating grease. Then, the device was kept in air for 1 h to keep the complete drying of the paste (Supplementary Fig. 18). For all electrical measurements, standard sample resistance ranging from 0.04 mΩ to 200 GΩ is within the detectable limitation, and the temperature range is from 5 to 400 K by helium cooling. After confirming the Ohm contact at different temperatures, we collected an I-V curve by scanning the current from 10 nA to 100 µA and measuring voltage at each step at every certain temperature. The electrical resistance was extracted from linear region of the I-V curve. The average resistivity ρ (Ω·m) of a sample is according to equation 1, where Rs (Ω·square) is the sheet resistance and t (m) is the thickness of sample. Hence the conductivity σ (S m-1) is calculated from Eq. 2.

$$\rho = R_s \cdot t$$
$$\sigma = \frac{1}{\rho } = \frac{1}{{R_s \cdot t}}$$

Time-resolved THz spectroscopy

The measurements were performed using an optical pump-THz probe setup driven by a titanium: sapphire laser amplifier system generating ~100 fs width laser pulses with a central wavelength of 800 nm, a pump fluence of 300 μJ cm-2, and a repetition rate of 1 kHz. The ~210-micron thick sample (O.D. (800 nm) = 2.3) was sandwiched between fused silica substrates and measured in transmission under nitrogen environment. The phase-sensitive detection of the THz pulse allows the evaluation of the complex conductivity of the photoexcited sample47. The frequency-resolved complex photoconductivity spectra of the sample were fitted by using the Drude-Smith (DS) model (Eq. 3):

$$\sigma _{{\mathrm{DS}}}\left( \omega \right) = \frac{{\omega _p^2\varepsilon _0\;\tau }}{{1 - i\;\omega \;\tau }} \ast (1 + \mathop {\sum}\limits_{n = 1}^\infty {\frac{{c_n}}{{\left( {1 - i\;\tau } \right)^n}}} )$$

where ωp, ε0, and are the plasma frequency, vacuum permittivity, scattering time and backscattering parameter, respectively. Mobility estimates on the DC limit were obtained from the formula obtained from the formula \(\sigma _{DS}\left( {\omega \to 0} \right) = {\mathrm{eN}}\mu (1 + {\mathrm{c}})\), where \(\mu = {\mathrm{e}}\frac{\tau }{{m^ \ast }}\).

Superconducting quantum interference device (SQUID)

Magnetometry was performed by using a SQUID-VSM (Quantum Design). Temperature dependence of the magnetization of the layered MOF powder sample was measured in zero-field cooling (ZFC) and field cooling (FC) sequence with applied magnetic field of 100 Oe. The magnetic field dependence of magnetization was measured at different temperatures, i.e., 5, 10, 20, 50, 100, 300, and 350 K.

57Fe Mössbauer measurements

Mössbauer measurements were carried out in a Cryo Vac helium flow cryostat with 6 litre helium volume protected by a nitrogen heat shield. We used commercial NIM rack devices. The drive was a Mössbauer WissEL drive unit MR-360 biased by a DFG-500 frequency generator in sinusoidal mode. We used for the data recording a CMTE multichannel data processor MCD 301/8 K and a WissEL single channel analyzer Timing SCA to select the energy window. The detector was a proportional counter tube and the source a Rh/Co source with an initial activity of 1.4 GB. The sample was a powder. The analysis was done by Moessfit48.

Modelling and electronic structure

Density functional theory (DFT) calculations were carried out using the Vienna ab-initio Simulation Package (VASP)49 version 5.4.1. The electronic wave-functions were expanded in a plane-wave basis set with a kinetic energy cutoff of 500 eV. Electron-ion interactions were described using the projector augmented wave (PAW) method50. Generalized gradient approximation (GGA)51 of the exchange-correlation energy in the form of Perdew-Burke-Ernzerhoff (PBE)52 was applied. We used DFT + U approach53 to describe the localized d-orbitals of Fe ions. The effective Coulomb (U) and exchange (J) terms were set to 4 and 1 eV, respectively. Such combination of U and J were already successfully applied for very similar systems28. The MOF monolayer was modeled by adding a large vacuum space of 10 Å in the direction normal to the monolayer. Monkhorst-Pack54 Gamma-centered grid with 6\(\times\)6\(\times\)1 dimension was used for k-point sampling of the Brillouin zone. In the computational protocol for the 3D stacking of the studied MOF the k-point grid dimension was changed to 6 × 6 × 5 and Grimme-D2 correction55 was applied.