A semiconducting layered metal-organic framework magnet

The realization of ferromagnetism in semiconductors is an attractive avenue for the development of spintronic applications. Here, we report a semiconducting layered metal-organic framework (MOF), namely K3Fe2[(2,3,9,10,16,17,23,24-octahydroxy phthalocyaninato)Fe] (K3Fe2[PcFe-O8]) with spontaneous magnetization. This layered MOF features in-plane full π-d conjugation and exhibits semiconducting behavior with a room temperature carrier mobility of 15 ± 2 cm2 V−1 s−1 as determined by time-resolved Terahertz spectroscopy. Magnetization experiments and 57Fe Mössbauer spectroscopy demonstrate the presence of long-range magnetic correlations in K3Fe2[PcFe-O8] arising from the magnetic coupling between iron centers via delocalized π electrons. The sample exhibits superparamagnetic features due to a distribution of crystal size and possesses magnetic hysteresis up to 350 K. Our work sets the stage for the development of spintronic materials exploiting magnetic MOF semiconductors.

S ince 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 device [1][2][3][4] . However, the practical realization of spintronic devices remains challenging due to the lack of functional materials displaying both, semiconducting and hightemperature magnetic ordering properties [5][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-VI 8 , III-V 9 , and IV-VI semiconductors 10,11 . Recently, intrinsic ferromagnetism has been demonstrated in layered inorganic two-dimensional (2D) materials, such as chromium germanium telluride (Cr 2 Ge 2 Te 6 ) 12 and chromium triiodide (CrI 3 ) 13 . Despite enabling long-range electronic coupling via short metal-metal bonds in these inorganic magnets 14 , the low magnetic ordering temperature (e.g., Tc = 159 K for GaMnAs, 61 K for Cr 2 Ge 2 Te 6 , and 68 K for CrI 3 , 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 magnets 15 . 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 MOFs 16 . Magnetic ordering at low temperature (Tc < 180 K, seen in Supplementary Table 2) has been already proven in bulk (3D) MOFs 14 . 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 conductivity 17 , due to the charge transfer of guests to the metal nodes or organic linkers. In addition, linking redox-active ligands, such as 2,5dihydroxybenzoquinone 14,18 , pyrazine 19 , 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 2012 20 , conjugated MOFs with layered structures have been developed displaying high electrical conductivities (up to 10 5 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 planes [20][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 temperature 28 . 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 K 3 Fe 2 [PcFe-O 8 ] MOF with square lattice geometry and interplane van der Waals (vdW) layer-stacking structure by employing (2,3,9,10,16,17,23,24- 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 K 3 Fe 2 [PcFe-O 8 ], 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 Synchrotron powder X-ray adsorption spectroscopy (XAS) was employed to further analysis the chemical state and coordination geometry in K 3 Fe 2 [PcFe-O 8 ]. Reference samples, e.g., Fe foil, FeO, Fe 2 O 3 as well as the precursor PcFe-OH 8 were also investigated by XAS. Fig. 1d shows the X-ray adsorption nearedge structure (XANES) spectra at K edge of all the samples. The Fe K-edge of XANES in K 3 Fe 2 [PcFe-O 8 ] exhibits a near-edge spectra similar to that of PcFe-OH 8 monomers, but completely different from those of FeO, Fe 2 O 3 , and Fe foil. Besides, the preedge feature (magnified in Fig. 1d) in K 3 Fe 2 [PcFe-O 8 ] 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 located 34 . Figure 1e displays the Fourier transform of the κweighted extended X-ray absorption fine structure (EXAFS) of K 3 Fe 2 [PcFe-O 8 ] as well as the contrast samples. The EXAFS presents a predominant peak in K 3 Fe 2 [PcFe-O 8 ], which is originated from the nearest-neighboring nitrogen or oxygen coordination shell around the Fe atoms 35 . 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  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 semiconductors 38 . 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 K 3 Fe 2 [PcFe-O 8 ] (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-O 4 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-O 4 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.
We further investigated the electronic band structure and PDOS of multi-layered K 3 Fe 2 [PcFe-O 8 ]. 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 K 3 Fe 2 [PcFe-O 8 ] 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 E ex = 300 meV (E ex = E AFM − E FM ), which implies the possible intrinsic ferromagnetic ordering of K 3 Fe 2 [PcFe-O 8 ] (Supplementary Fig. 17).
The electrical conductivity of the K 3 Fe 2 [PcFe-O 8 ] samples through van der Pauw method was derived to be 2 × 10 − 3 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 K 3 Fe 2 [PcFe-O 8 ] 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 K 3 Fe 2 [PcFe-O 8 ] 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 energy 23,40 . In addition, it needs to mention that redox state of the iron centers also contributes to the conductivity of the K 3 Fe 2 [PcFe-O 8 ]. 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 Fe 2+ into Fe 3+ with a ratio of~1/3 in the oxidized K 3 Fe 2 [PcFe-O 8 ] (Supplementary Fig. 21b).
The Hall resistance (R Hall ) was further measured under magnetic field (H) based on the van der Pauw pattern ( Supplementary Fig. 22), which displays a linear increase of R Hall with H. The fitting slope of the Hall effect shows a p-type   Supplementary Fig. 23). Next, contact-free time-resolved THz spectroscopy (TRTS) was employed to further address the photoconductivity and mobility nature in the K 3 Fe 2 [PcFe-O 8 ] samples. Figure 3b shows the real part of the photoconductivity measured as a function of pumpprobe 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*~1.9m 0 (obtained from DFT calculations), we obtain a high mobility estimate for the photo-generated free carriers at the peak conductivity of 15 ± 2 cm 2 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 scattering 10 , indicating the interaction between the magnetic moments and the carriers.

Magnetic properties of the K 3 Fe 2 [PcFe-O 8 ].
To probe the magnetic behavior of the K 3 Fe 2 [PcFe-O 8 ], fieldand 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, K 3 Fe 2 [PcFe-O 8 ] 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 μ B 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 K 3 Fe 2 [PcFe-O 8 ] 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). 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 K 3 Fe 2 [PcFe-O 8 ] (shown in the TEM images, Supplementary Figs. 10 and 11). The nanoscale K 3 Fe 2 [PcFe-O 8 ] 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 πelectrons along the ab plane contributes to the persistence of spontaneous magnetization above room temperature. Compared with the Curie temperatures of thus far reported magnetic MOFs 14,18,42-44 , 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.

Discussion
To further understand the magnetic mechanism in K 3 Fe 2 [PcFe-O 8 ], we investigated the local environment of Fe atoms by 56 Fe 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 ΔE Q = 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 57 Fe Mössbauer spectrum demonstrates the only presence of one Fe (III) state in the monomer PcFe-OH 8 (Supplementary Fig. 26), which reveals that phthalocyanine ligands contribute to partial Fe site 1 while the Fe site 2 originates from the linkages in K 3 Fe 2 [PcFe-O 8 ] MOFs. This observation is consistent with the ferromagnetic exchange within the individual clusters 45 . Upon cooling from room temperature to 25 K, three Fe sites are observed in K 3 Fe 2 [PcFe-O 8 ] (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 ; ΔE Q = 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 56 Fe 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 K 3 Fe 2 [PcFe-O 8 ] 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 backbones 10,24 , thus enabling strong hybridization between the d/p orbitals of Fe, the phthalocyanine core, and the Fe-O 4 nodes (seen in Fig. 2c). As mentioned before, the calculated magnetic ground states of AA-serrated K 3 Fe 2 [PcFe-O 8 ] 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 (K 3 Fe 2 [PcFe-O 8 ]) with layered structures, which exhibits a p-type semiconducting behavior with a high mobility of~15 cm 2 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.

Methods
Materials. 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.
All the reactions were performed under vacuum using the Schlenk line technique. 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 R s (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.
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 sample 47 . The frequency-resolved complex photoconductivity spectra of the sample were fitted by using the Drude-Smith (DS) model (Eq. 3): where ω p , ε 0, 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 σ DS ω ! 0 ð Þ¼eNμð1 þ cÞ, where μ ¼ e τ m Ã .
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 zerofield 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.
57 Fe 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 Moessfit 48 .
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) method 50 . 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 approach 53 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 systems 28 . The MOF monolayer was modeled by adding a large vacuum space of 10 Å in the direction normal to the monolayer. Monkhorst-Pack 54 Gamma-centered grid with 6 6 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 correction 55 was applied.

Data availability
The data that support the findings of this study are available from the corresponding author on reasonable request.