Sujing Wang

Nertil Xhaferaj

Mohammad Wahiduzzaman

Kolade Oyekan

, Xiao Li

Kevin Wei

Bin Zheng

Antoine Tissot

Jérôme Marrot

William Shepard

Charlotte Martineau-Corcos

Yaroslav Filinchuk

Kui Tan
email: kuitan@utdallas.edu

Guillaume Maurin
email: guillaume.maurin1@umontpellier.fr

Christian Serre
email: christian.serre@ens.fr

Engineering Structural Dynamics of Zirconium Metal-Organic Frameworks Based on Natural C4-linkers

Engineering the structural flexibility of MOF materials for separation-related applications remains a great challenge. We present here a strategy of mixing rigid and soft linkers in a MOF structure to achieve tunable structural flexibility, as exemplified in a series of stable isostructural Zr-MOFs built with natural C4 linkers (fumaric acid, succinic acid and malic acid). As shown by the differences in linker bond stretching and rotational freedom, these MOFs display distinct responsive dynamics to external stimuli, namely temperature or guest adsorption. Comprehensive in-situ characterizations reveal a clear correlation between linker character and MOF dynamic behavior, which leads to the discovery of a multivariate flexible MOF. It shows an optimal combination of both good CO 2 working capacity and significantly enhanced CO 2 /N 2 selectivity . In principle, it provides a new avenue for potentially improving the ability of microporous MOFs to separate other gaseous and liquid mixtures.

INTRODUCTION

Flexible metal-organic frameworks (MOFs) or porous coordination polymers (PCPs) represent a particular subgroup within the class of the hybrid materials, namely the third generation compounds or soft porous crystals (SPCs). [START_REF] Kitagawa | Functional micropore chemistry of crystalline metal complex-assembled compounds[END_REF][START_REF] Bureekaew | Chemistry and application of flexible porous coordination polymers*[END_REF] In contrast to the first and second generations, [START_REF] Liu | Applications of metal-organic frameworks in heterogeneous supramolecular catalysis[END_REF][START_REF] Horike | Soft porous crystals[END_REF][START_REF] Kitagawa | Functional porous coordination polymers[END_REF] these materials display intriguing structural dynamics in response to external stimuli, such as guest molecule variations, [START_REF] Seo | A pillared-layer coordination polymer with a rotatable pillar acting as a molecular gate for guest molecules[END_REF][START_REF] Loiseau | A Rationale for the Large Breathing of the Porous Aluminum Terephthalate (MIL-53) Upon Hydration[END_REF][START_REF] Mellot-Draznieks | Very large swelling in hybrid frameworks: a combined computational and powder diffraction study[END_REF][START_REF] Krause | A Stimuli-Responsive Zirconium Metal-Organic Framework Based on Supermolecular Design[END_REF][START_REF] Qin | Flexible Zirconium MOF as the Crystalline Sponge for Coordinative Alignment of Dicarboxylates[END_REF][START_REF] Zhang | A Flexible Metal-Organic Framework with 4-Connected Zr6 Nodes[END_REF] light, [START_REF] Brown | Photophysical pore control in an azobenzene-containing metal-organic framework[END_REF][START_REF] Modrow | The first porous MOF with photoswitchable linker molecules[END_REF] temperature, [START_REF] Liu | Reversible structural transition in MIL-53 with large temperature hysteresis[END_REF][START_REF] Devries | Nanoscale lattice fence" in a metal-organic framework: interplay between hinged topology and highly anisotropic thermal response[END_REF][START_REF] Henke | Massive Anisotropic Thermal Expansion and Thermo-Responsive Breathing in Metal-Organic Frameworks Modulated by Linker Functionalization[END_REF] electrical field [START_REF] Ghoufi | Electrically Induced Breathing of the MIL-53(Cr) Metal-Organic Framework[END_REF] and pressure. [START_REF] Chapman | Pressure-induced amorphization and porosity modification in a metal-organic framework[END_REF][START_REF] Gagnon | MOFs under pressure: the reversible compression of a single crystal[END_REF][START_REF] Goodwin | Large negative linear compressibility of Ag3[Co(CN)6][END_REF][START_REF] Yot | Large breathing of the MOF MIL-47(VIV) under mechanical pressure: a joint experimental-modelling exploration[END_REF] Consequently, numerous unique and unprecedented attributes have been discovered with regard to their applications in areas such as separation, sensors, catalysis, and biomedicine. [START_REF] Schneemann | Flexible metal-organic frameworks[END_REF][START_REF] Lin | Metal-organic frameworks based on flexible ligands (FL-MOFs): structures and applications[END_REF][START_REF] Férey | Giant flexibility of crystallized organic-inorganic porous solids: facts, reasons, effects and applications[END_REF][START_REF] Férey | Structural flexibility in crystallized matter: from history to applications[END_REF] In this context, controlled inducement of the structural flexibility during MOF synthesis was regarded as an attractive strategy for targeting a given application. However, successful implementation remains a great challenge.

There are four major categories of flexibility identified for MOFs in the literature, including breathing, swelling, linker rotation and sub-network displacement. [START_REF] Bousquet | Adsorption induced transitions in soft porous crystals: an osmotic potential approach to multistability and intermediate structures[END_REF][START_REF] Coudert | Adsorption Deformation and Structural Transitions in Metal-Organic Frameworks: From the Unit Cell to the Crystal[END_REF][START_REF] Schneemann | Flexible metal-organic frameworks[END_REF] In the case of breathing MOFs, such as the MIL-53 (Materiaux from Institute Lavoisier) series, the influence of phenylene ring rotation about its C-C axis to different extends are responsible for the notable changes in the unit cell. [START_REF] Liu | Reversible structural transition in MIL-53 with large temperature hysteresis[END_REF] A recent discovery of a breathing MOF displaying negative gas adsorption properties, namely DUT-49 (Dresden University of Technology No. 49), [START_REF] Krause | A pressure-amplifying framework material with negative gas adsorption transitions[END_REF] showed that metal node rotation could also lead to structural breathing, with a noticeable induced-bending of the biphenyl moiety in the linker . For the swelling examples, such as MIL-88, metal node rotation is the sole driving force, with a very limited contribution of the linker to the overall change in the unit cell parameters . [START_REF] Férey | Large breathing effects in three-dimensional porous hybrid matter: facts, analyses, rules and consequences[END_REF] Finally, linker rotation that does not induce significant changes in the unit cell has been observed in some non-breathing/swelling MOF structures such as ZIF-8 and MIL-140s, which are built with linkers containing rigid aromatic centers. [START_REF] Schneemann | Flexible metal-organic frameworks[END_REF] To the best of our knowledge, each flexible MOF reported has its individual combination of linker selection and inorganic building block. Consequently, the corresponding structural flexibility can only be tuned by the introduction of attached functional groups. The resulting linker bond distortion and bending are thus unrelated to the local freedom of bond stretching and rotation. On the other hand, flexible aliphatic linkers with linear or cyclic alkane centers that possess greater bond stretching and rotational freedom were expected to favor the generation of dynamic structures. However, relevant studies in this field remain scarce. [START_REF] Lin | Metal-organic frameworks based on flexible ligands (FL-MOFs): structures and applications[END_REF][START_REF] Katsoulidis | Chemical control of structure and guest uptake by a conformationally mobile porous material[END_REF][START_REF] Bueken | A Flexible Photoactive Titanium Metal-Organic Framework Based on a [Ti(IV)3(mu3-O)(O)2(COO)6] Cluster[END_REF][START_REF] Bueken | A Breathing Zirconium Metal-Organic Framework with Reversible Loss of Crystallinity by Correlated Nanodomain Formation[END_REF] . Hence, it is very difficult to investigate the impact of linker ' freedom of bond stretching and rotation on the softness of a given MOF type when the rigid and flexible linker pair is unable to generate isostructural MOFs. In this regard, it would be of great interest to engineer the structural flexibility in a series of isostructural MOFs using the same inorganic building block, and linkers of similar molecular size and configuration but with different degrees of structural freedom (in both bond stretching and rotation), in order to understand how the local and structural dynamics of the architectures are governed by the linker characters.

To achieve this goal, the expected MOF candidates should fulfill the following requirements: 1) isostructural analogues built with linkers of different degrees of freedom; 2) an inorganic building block exhibiting a possible plane of symmetry, allowing the cooperative movements of various bonds involved; 3) frameworks featuring an accessible porosity to certain stimuli that can easily be detected, analyzed and compared; 4) compounds being robust to avoid any structural deterioration.

RESULTS AND DISCUSSION

Among the reported inorganic building blocks having a possible symmetry plane, the Zr 6 cluster is one of the best options owing to its high connectivity and chemical stability, natural abundance, and low toxicity. The Zr 6 O 4 (OH) 4 cluster is one of the most common secondary building units (SBUs) reported in the literature. [START_REF] Bai | Zr-based metal-organic frameworks: design, synthesis, structure, and applications[END_REF][START_REF] Gu | Design and control of gas diffusion process in a nanoporous soft crystal[END_REF] The varying node connectivity, including the 12-, 10-, 8-, 6-, and 4-connected Zr 6 -oxocluster in MOFs, allow a high degree of tunability of the corresponding MOFs structures, which provides a better opportunity to achieve a flexible crystal architecture. The 12-connected Zr 6 cluster has the densest connection environment and thus less freedom and space for configuration rearrangement. Alternatively, reducing the connectivity of the Zr 6 cluster could be more favorable to introduce structural flexibility. [START_REF] Zhang | A Flexible Metal-Organic Framework with 4-Connected Zr6 Nodes[END_REF] In view of these factors, Zr-MOFs with low connectivity SBUs were preferentially considered as the model compounds.

Judicious use of flexible linkers has already been demonstrated to be an efficient strategy for generating soft MOFs structures when the inorganic SBUs are rigid. [START_REF] Lin | Metal-organic frameworks based on flexible ligands (FL-MOFs): structures and applications[END_REF][START_REF] Schneemann | Flexible metal-organic frameworks[END_REF] Among the reported flexible linkers, naturally occurring compounds are of particular interest since they are bio-compatible and thus efficiently reduce the toxicity of the resulting MOFs. In the limited cases of existing Zr-MOFs composed of natural acids, [START_REF] Bai | Zr-based metal-organic frameworks: design, synthesis, structure, and applications[END_REF] fumaric acid and its derivatives are dominant despite the fact that they all share the 12-fold connectivity, which highlights the controllable crystallization of Zr-fumarate type frameworks, [START_REF] Furukawa | Water adsorption in porous metal-organic frameworks and related materials[END_REF][START_REF] Sk | Synthesis, Characterization, Stability, and Gas Adsorption Characteristics of a Highly Stable Zirconium Mesaconate Framework Material[END_REF][START_REF] Wang | A Microporous Zirconium Metal-Organic Framework Based on Trans-aconitic Acid for Selective Carbon Dioxide Adsorption[END_REF][START_REF] Wang | A robust zirconium amino acid metal-organic framework for proton conduction[END_REF] as well the rigid structural character of fumaric acid, which is the same as evidenced in other MOFs. [START_REF] Férey | Giant flexibility of crystallized organic-inorganic porous solids: facts, reasons, effects and applications[END_REF][START_REF] Férey | Structural flexibility in crystallized matter: from history to applications[END_REF][START_REF] Elsa | The Structure of the Aluminum Fumarate Metal-Organic Framework A520[END_REF] On the contrary, succinic acid, which behaves as a soft molecular spring as seen in Co-MOFs, [START_REF] Livage | Hybrid open frameworks (MIL-n). Part 5 -Synthesis and crystal structure of MIL-9: a new three-dimensional ferrimagnetic cobalt(II) carboxylate with a two-dimensional array of edge-sharing Co octahedra with 12-membered rings[END_REF][START_REF] Livage | Hybrid Open Networks (MIL 16): Synthesis, Crystal Structure, and Ferrimagnetism of Co4(OH)2(H2O)2(C4H4O4)3•2H2O, a New Layered Cobalt(II) Carboxylate with 14-Membered Ring Channels[END_REF] has molecular lengths and configurations similar to those of fumaric acid. The most noticeable difference between succinic acid and the double-b f m f m ' single-bonded carbon skeleton, which usually exhibits greater freedom of stretching and rotation. In view of these features, fumaric and succinic acids, as the simplest pair of rigid and flexible natural carboxylic linkers respectively, were selected for their typicality and suitability for the aforementioned target. Additionally, malic acid, a succinic acid derivative with a hydroxyl group on the alkane chain, was also used to check the impact of the side functional group on the flexibility of the architecture.

Fumaric acid (Fig. 1a) was first used to react with ZrCl 4 as model reaction. Formic acid has been proven to be a modulating agent to generate the 12-connected structure with improved crystallinity in the preparation of Zr-Fumarate (as known as MOF-801). [START_REF] Wißmann | Modulated synthesis of Zr-fumarate MOF[END_REF] We expected that the increase in the amount of formic acid in the reaction solution could decrease the connectivity of the Zr 6 cluster, leading to less rigid structures with lower structural symmetry. Pure formic acid was used as solvent in this case since it has been shown to be an efficient modulator for preparing group IV metal-based MOFs. [START_REF] Wang | A Microporous Zirconium Metal-Organic Framework Based on Trans-aconitic Acid for Selective Carbon Dioxide Adsorption[END_REF][START_REF] Wang | A phase transformable ultrastable titanium-carboxylate framework for photoconduction[END_REF] As expected, instead of forming the well-known Zr-Fumarate structure, MIP-203-F, a new phase with lower structural symmetry was obtained (Fig. 1b). When succinic acid was used under the same reaction conditions, a clear solution was observed without any solid product, possibly due to the much greater solubility of succinic acid in formic acid. After optimizing the reactant concentrations and the ratio between ZrCl 4 and succinic acid, the expected isostructural product was successfully obtained. The synthesis of MIP-203-M could be achieved by tuning the reactant ratio under similar conditions (see SI for detail). The crystal structure model of MIP-203 was initially determined on MIP-203-S using a dual computational and experimental approach developed on the principle of Automated Assembly of Structure Building Unit (AASBU) theory that combines high-resolution powder X-ray diffraction data and our in-house crystal structure prediction software. [START_REF] Wahiduzzaman | Computational structure determination of novel metal-organic frameworks[END_REF] Afterwards, needle-like single crystals of MIP-203-F with a homogeneous size and shape distribution (Fig. 1c and1d) were subjected to the synchrotron diffraction single crystal data collection with a micro-focused X-rays on the Proxima 2A beamline (Synchrotron SOLEIL, France), [START_REF] Duran | PROXIMA 2A -A New Fully Tunable Micro-focus Beamline for Macromolecular Crystallography[END_REF] in order to experimentally confirm the major framework connectivity of the simulated structural model.

Since the three MIP-203 compounds are isostructural, MIP-203-F was used as an example to detail their structural features. It was found that MIP-203-F,

Zr 6 (μ 3 -O) 4 (μ 3 -OH) 4 (fumarate) 4 (formate) 2 (OH) 2 (H 2 O) 2 ,

crystallizes in an orthorhombic

Immm space group with unit cell parameters of a = 10.000(2) Å, b = 11.940(2) Å, c = 19.829 (11) Å and V = 2367.6(15) Å [START_REF] Liu | Applications of metal-organic frameworks in heterogeneous supramolecular catalysis[END_REF] . The Zr 6 SBUs (Fig. 1e) were interconnected with one another by both fumarate and formate linkers, resulting in the final three-dimensional (3D) framework. It is, to the best of our knowledge, the first time that a Zr 6 SBU with mixed linkage from two highly dissimilar linkers in term of shape, configuration and connectivity was generated using direct synthesis for Zr-MOFs. Formate serves as the auxiliary ligand connecting the adjacent SBUs along the c-axis. However, their presence blocks the window of the pores along the a-axis (Fig. 1f). When the structure is viewed along the b-axis (Fig. 1g), formates are shown to occupy the middle of the large rhombic channels, dividing them into two 4 Å sized triangular channels, which is also observed in the case of UiO-66(Zr) type structures. [START_REF] Cavka | A new zirconium inorganic building brick forming metal organic frameworks with exceptional stability[END_REF] Due to the presence of the terminal OH -/H 2 O pairs in the SBU, there exist rhombic channels running along the c-axis without any blocking species (Fig. 1h).

Simulation results indicate that both MIP-203-F and MIP-203-S possess almost identical pore volumes (0.27 and 0.25 cm 3 g -1 for -F and -S, respectively) and pore sizes estimated from the PSD plots (largest pore diameters ~5 Å, Figure S3).

However, due to the difference between C-C single and C=C double bonds, fumaric and succinic acids show slightly different conformations (see Figure S4-S6), resulting in the pore shape and dimension of the corresponding MOFs to vary to some extent (Figure S3). The cell volume of the activated samples decreased by 57.6 Å 3 /2.4%, 73.9 Å 3 /3.1% and 80.4 Å 3 /3.4% for MIP-203-F, -M and -S, respectively, in comparison with those of the as-synthesized ones, suggesting their similar structural contraction abilities upon thermal stimulus (Table S2, S3 and S4).

In situ Raman spectroscopy was further employed to monitor the changes occurring in specific bonds in the MOF structure upon activation. Figs 3a and 3b present the most pronounced changes in the range of 1800-1200 cm -1 caused by in situ activation of the sample at 100 °C (see full spectra in Supporting information). In comparison to the as-synthesized MIP-203-S, the peak centers of the carboxylate stretching bands ν as (COO) ν s (COO) in the activated sample are blue-shifted of 7 cm -1 and 6 cm -1 , respectively, indicative of the bond hardening due to structural contraction upon removal of the solvents. T p g ν s (COO) change in MIP-203-F sample is only 3 cm -1 ν as (COO) peak intensity is too weak to be addressed, which evidences a much less pronounced bond change compared to that of MIP-203-S. Furthermore, a band at 1289 cm -1 for activated MIP-203-S with a 11 cm -1 red-shift compared to the as-synthesized one was assigned to the vibration mode that involves -CH 2 wagging coupled with C-C-C-C stretching (see the attached videos 1 and 2). On the contrary, only 7 cm -1 difference (1277 to 1284 cm -1 ) for the -CH wagging in C-C=C-C was detected for MIP-203-F before and after the thermal activation. This mode has been shown to be very sensitive to changes in the bond angle of the carbon chain. [START_REF] Nijem | Tuning the gate opening pressure of Metal-Organic Frameworks (MOFs) for the selective separation of hydrocarbons[END_REF] As shown in Figures 3a and3b by DFT calculation, when the angles p and q (p and q are bond angles of the C4 skeleton of the linker) are decreased by ~2° and ~4° respectively for fumaric acid, and ~5° and ~5° for succinic acid, its frequency shifts downward by ~10 cm -1 and ~15 cm -1 in fumaric acid and succinic acid, respectively.

The trend of these shifts is in good agreement with the experimentally observed values. It also matches well the similar observation reported on other flexible MOFs. [START_REF] Nijem | Tuning the gate opening pressure of Metal-Organic Frameworks (MOFs) for the selective separation of hydrocarbons[END_REF] The Brunauer-Emmett-Teller (BET) area of 430 m 2 g -1 and a total pore volume of 0.24 cm 3 g -1 . In a sharp contrast, MIP-203-S did not show any accessible porosity to N 2 at 77 K (Fig. 4a). It is likely that soft single-bonded C4 skeleton of succinic acid that can distort more than fumaric acid with C=C double bond. After thermal activation under vacuum to remove the guest molecules, MIP-203-S tends to stay in a closely packed form so that it is not accessible to nitrogen. When the adsorption temperature was increased to 273 K, the thermodynamic motions of single-bonded aliphatic chain of succinate resulted in a slight increase of N 2 uptake (Figure S12). The much higher uptake for MIP-203-F at 273 K suggests its more pronounced swelling ability towards nitrogen sorption stimuli or larger pore opening in its activated form. In the case of MIP-203-M, it shows a BET area of 380 m 2 g -1 and pore volume of 0.21 cm When the probe molecule was switched to the smaller-sized CO 2 , the influence of linker softness on the structural dynamics is much less pronounced. MIP-203-S exhibits half the uptake of CO 2 compared to MIP-203-F over almost the entire pressure range at 273 K (Figure S13). However, a significant decrease in CO 2 uptake was observed for MIP-203-S when the adsorption temperature was increased to 298 K.

The uptake at 1 bar drops by 58%, from 1.61 mmol g -1 at 273 K to 0.68 mmol g -1 at 298 K. In the case of MIP-203-F, only a 24% reduction (2.72 mmol g -1 to 2.05 mmol Furthermore, Ideal adsorbed solution theory (IAST) [START_REF] Bae | Carborane-based metal-organic frameworks as highly selective sorbents for CO(2) over methane[END_REF][START_REF] Heuchel | Adsorption of CH4-CF4 Mixtures in Silicalite: Simulation, Experiment, and Theory[END_REF] was applied to evaluate the adsorption CO 2 /N 2 selectivity of MIP-203s when a binary mixture of CO 2 :N 2 =15:85 (v:v) at 298 K was considered (Fig. S14). The corresponding selectivities were

  Herein, we present a series of isostructural MOFs based on 10-connected Zr 6 clusters and natural C4 carboxylate linkers with linear aliphatic chains, denoted as MIP-203-F, MIP-203-S, and MIP-203-M (MIP stands for the Materials of the Institute of Porous Materials from Paris, F for fumaric acid, S for succinic acid, and M for malic acid), designed to engineer structural flexibility based on the nature and degree of freedom of the linker. While MIP-203-F is constructed using double-bonded fumarate and MIP-203-S is built with single-bonded succinate, MIP-203-M contains a sparse framework displaying multivariate flexibility

Figure 1 .

 1 Figure 1. Structure details of MIP-203. (a) Chemical structures of the three naturally occurring dicarboxylic acids

An

  in situ dehydration reaction occurred when malic acid was introduced in the synthesis of MIP-203-M, leading to the transformation of 2/3 malic acid into fumaric acid. MIP-203-M can therefore be considered as the corresponding product of replacing 1/3 fumarate in MIP-203-F by malate, as evidenced by solid-state NMR data (Figure S7). The coexistence of fumarate and malate in MIP-203-M makes it an unusual example of the MOF framework that shows multivariate flexibility through direct synthesis. Thermal and chemical stability tests of MIP-203s were carried out before investigating and comparing the response to different external stimuli of the MIP-203s. The influence of thermal activation of the as-synthesized samples was first studied. As shown in Fig 2a, the as-synthesized MIP-203-F sample displays four PXRD peaks at low angle, i.e. 8.46, 8.94, 9.86 and 11.36° corresponding to the (011), (002), (101) and (110) Bragg peaks, respectively. After activating the sample at 100 °C under vacuum, the distinct shifts to higher angle domains for the peaks at 8.73 and 11.56° were observed., which could be ascribed to the notable decrease of the b unit-cell parameter in the activated sample (11.856 Å) in comparison with that of the as-synthesized one (12.246 Å). The larger change in b compared to a and c could be explained by the weaker connection along b-axis in the crystal structure while the presence of bridging formates undoubtedly enhances the overall rigidity of the a-c plane. Similar observation was noticed between the as-synthesized and activated samples of both MIP-203-S and MIP-203-M (Fig 2b and 2c).

Figure 2 .

 2 Figure 2. Structural response at room temperature of MIP-203s to liquid external stimuli (PXRD λ Cu ≈ 1,5406 Å). a. MIP-203-F in different solvents. b. MIP-203-S in different solvents. c. MIP-203-M in different solvents.

1 Figure 3 .

 13 Figure 3. Raman spectra of MIP-203-F (a) and MIP-203-S (b) in as-synthesized (top lines) and dried (bottom lines) form upon in situ activation. The spectra are collected at room temperature. In situ IR spectra of adsorbed DMF in MIP-203-F (c), MIP-203-S (d) and MIP-203-M (e), referenced to the activated MOFs (bottom black lines) under vacuum.

3 g - 1 comparable

 1 to that of MIP-203-F, indicating the beneficial effect of the fumarate content. A notable reduction of N 2 uptake was however observed at 273 K in comparison with that of MIP-203-F, showing the critical role of the soft malate content.

Figure 4 .

 4 Figure 4. Single component gases adsorption behaviors of MIP-203s. a N 2 adsorption isotherms collected at 77 K. b CO 2 and N 2 adsorption isotherms collected at 298 K.

g - 1 )

 1 in CO 2 uptake at 1 bar was observed from 273 K to 298 K. Therefore, MIP-203-F and MIP-203-S indeed display huge differences in their structural responses to gaseous guest molecules, as expected. MIP-203-M behaves more similarly to MIP-203-F than MIP-203-S, likely due to the major linker component of fumarate in its structural framework. To check the structural response of the MOF framework in the presence of CO 2 guest molecules, in situ PXRD data along the adsorption of CO 2 under various pressures was collected on the three MIP-203 samples at room temperature. As shown in Figure 5, both activated MIP-203-F and MIP-203-M structures could easily swell back to the open form under a low CO 2 pressure (between 200-300 mbar), highlighting their swelling flexibility. Further increasing the CO 2 pressure did not have any notable effects on the mp ' structural expansion. However, the activated MIP-203-S sample remained in its closely packed configuration throughout the entire range of CO 2 pressures applied. Even high CO 2 pressure (6 bar) was not able to reopen the framework, which is consistent with its lower CO 2 uptake compared to those of MIP-203-F and MIP-203-M under the same condition. To evaluate the adsorption affinity, isosteric heats of adsorption (Q st ) were calculated with the CO 2 adsorption isotherms collected at 273 K, 298 K and 303 K in the pressure range of 0-1 bar. At zero coverage, the Qst values are 32 KJ mol -1 , 34 KJ mol -1 and 33 KJ mol -1 for MIP-203-F, MIP-203-S and MIP-203-M, respectively. The working capacities of CO 2 sorption (Wc(1-0.1 bar)) for three MIP-203 samples were calculated to be 1.38, 1.04 and 1.15 mmol/g for MIP-203-F, MIP-203-S and MIP-203-M, respectively, at 273 K. The corresponding values at 298 K are 1.35, 0.47 and 1.09 mmol/g. It is interesting to note that the working capacities for both MIP-203-F and MIP-203-M showed far less dependence on the considered temperature, while the working capacity of MIP-203-S exhibited dramatic sensitivity towards the applied temperature, decreasing by more than half when the temperature increased from 273 K to 298 K.

Figure 5 .

 5 Figure 5. In situ PXRD patterns of MIP-203s along adsorption of CO 2 under different pressures at room temperature. a MIP-203-F. b MIP-203-S. c MIP-203-M.

calculated to be 15 ,

 15 34 and 51 for MIP-203-F, MIP-203-S and MIP-203-M, respectively (FigureS14), which are in the usual range for MOF-based CO 2 sorbents.[START_REF] Benoit | MIL-91(Ti), a small pore metal-organic framework which fulfils several criteria: an upscaled green synthesis, excellent water stability, high CO2 selectivity and fast CO2 transport[END_REF][START_REF] Liu | Progress in adsorption-based CO2 capture by metal-organic frameworks[END_REF] It is worth noting that MIP-203-S shows a selectivity twice as high as that of MIP-203-F, highlighting the linker softness effect that interferes nitrogen from entering the pores. However, MIP-203-M outperforms MIP-203-S in terms of selectivity by a margin of 50%. The above comparison of the CO 2 working capacity and adsorption selectivity between the three MIP-203 compounds suggests that the introduction of a proper local softness into a structurally flexible MOF framework would result in a favorable combination of accessible porosity and structural dynamics, leading to an optimal balance for selective CO 2 separation in terms of working capacity and selectivity.CONCLUSIONIn summary, we report a series of flexible isostructural Zr-MOFs synthesized based on natural C4 linkers (MIP-203s), in which the structural dynamics can be engineered based on the differences in freedom of linker bond stretching and rotation between rigid fumarate and soft succinate. Comprehensive characterizations of structural responses to diverse external stimuli, including thermal changes as well as liquid and gaseous guest molecules, have been conducted to correlate the corresponding linker' flexibility with MOF dynamics. Directed by this established correlation, MIP-203-M, with a proper multivariate flexibility that efficiently combines swelling and linker distortion/bending, was evaluated for the adsorptive separation of CO 2 /N 2 . It not only exhibits a good working capacity for CO 2 adsorption under the working condition benefiting from the swelling of MIP-203-F framework, but also a significantly enhanced selectivity for CO 2 /N 2 compared to those of the -F and -S forms, owing to the linker stretching and rotational flexibility generated from the soft succinate skeleton. In principle, this strategy could be applied to the separation of other chemical mixtures, including liquid and gaseous phases. Therefore, it serves as a suitable example that engineering the structural flexibility of MOF materials could constitute a promising and versatile strategy for improving MOF separation performance, or even realizing new applications.

ACKNOWLEDGEMENT

The authors from France are grateful to the ANR Project MeaCoPA (ANR-17-CE29-0003) for financial support. The spectroscopic characterization and analysis (IR and Raman) were supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences under Award Number DE-FG02-08ER46491

AUTHOR INFORMATION

Notes

The authors declare no competing interest.


Mohamed Helmi

Hadj Alouane
email: malouane@kfu.edu.sa

Olfa Nasr

Hammadi Khmissi

Bouraoui Ilahi

Gilles Patriarche

Mohamad Mahmoud Ahmad

Michel Gendry

Catherine Bru-Chevallier

Nicolas Chauvin

M H Hadj Alouane

Temperature dependence of optical properties of InAs/InP quantum rod-nanowires grown on Si substrate

Keywords:

III-V nanowires

Introduction

III-V nanowires (NWs), such as InAs, GaAs or InP, are ideal for photonic devices due to their direct band gap and superior electrical properties [START_REF] Joyce | Phase perfection in zinc blende and wurtzite III-V nanowires using basic growth parameters[END_REF]. These NWs have been proposed for a wide variety of applications ranging from interconnects to functional device elements in electronic and optoelectronic applications [START_REF] Sumikura | Mid-infrared lasing of single wurtzite InAs nanowire[END_REF][START_REF] Gibson | Journal of Luminescence xxx (xxxx) xxx-xxx broadband high-speed single-photon detection[END_REF][START_REF] Barrigón | GaAs nanowire pn-junctions produced by low-cost and high-throughput aerotaxy[END_REF][START_REF] Bai | InAs/InP core/shell nanowire gas sensor: effects of InP shell on sensitivity and long-term stability[END_REF][START_REF] Hill | Interplay of surface recombination and diode geometry for the performance of axial p-i-n nanowire solar cells[END_REF]. Such structures are typically fabricated by a bottom-up approach using the Vapor-Liquid-Solid (VLS) method [START_REF] Wagner | Vapor-liquid-solid mechanism of single crystal growth[END_REF]. The NW growth using VLS method is well adapted to grow III-V semiconductors on silicon (Si) substrate which enables III-V NW optoelectronic devices to be integrated with the established Si microelectronics technology [START_REF] Woo | Effect of twinning on the photoluminescence and photoelectrochemical properties of indium phosphide nanowires grown on silicon (111)[END_REF]. In contrast to the stable zinc blende (ZB) phase in bulk III-V materials, the Au catalyst growth of III-V NWs occurs preferentially in the hexagonal wurtzite (WZ) phase [START_REF] Tchernycheva | Growth and characterization of InP nanowires with InAsP insertions[END_REF][START_REF] Mattila | Crystal-structure-dependent photoluminescence from InP nanowires[END_REF][START_REF] Minot | Single quantum dot nanowire LEDs[END_REF] in which the valence band splits in 3 bands, the so called A, B and C bands [START_REF] Gadret | Valence-band splitting energies in wurtzite InP nanowires: photoluminescence spectroscopy and ab initio calculations[END_REF]. Furthermore, NWs with built-in homojunctions or heterojunctions, realized by variation of doping and composition along their length, constitute a new class of promising one dimensional semi conductor heterostructures [START_REF] Gudiksen | Growth of nanowire superlattice structures for nanoscale photonics and electronics[END_REF][START_REF] Panev | Sharp exciton emission from single InAs quantum dots in GaAs nanowires[END_REF]. The insertion of low dimensional quantum heterostructures in NWs adds new degrees of design freedom for optimization of electrical and optical performances [START_REF] Karimi | Room-temperature InP/InAsP quantum discs-in-nanowire infrared photodetectors[END_REF]. Generally, the VLS approach allows for the growth of quantum rods of any desired height without the limits inherent to the growth of InAs quantum dots using the Stranski-Krastanov (SK) growth mode [START_REF] Zieliński | Fine structure of light-hole excitons in nanowire quantum dots[END_REF]. Accordingly, InP NWs with InAs QRs are particularly interesting for fiber optic communication applications through the ability to reach the 1.55 μm emission wavelength at room temperature [START_REF] Tchernycheva | Growth and characterization of InP nanowires with InAsP insertions[END_REF]. The difference on the crystal lattice between the InAs QR and the InP NW produces an elastic strain which affects the optoelectronic properties of such nanostructures [START_REF] Ferrand | Strain in crystalline core-shell nanowires[END_REF]. To ensure efficient optoelectronic devices operating at room temperature, a clear understanding of the effect of temperature on the optical properties of such structure is also highly required. In this context, we report on the optoelectronic properties of InAs/InP quantum rod-nanowires (QR-NWs) grown on Si substrates using the finite element method (FEM) taking into account the strain effect and we investigate their temperature dependent photoluminescence (TDPL) properties.

Materials and methods

The InAs/InP QR-NWs are grown on Si(001) substrates by solid-source molecular beam epitaxy (ssMBE) and VLS method using Au-In droplets obtained by dewetting of a Au-In film at 700 °C [START_REF] Anufriev | Polarization properties of single and ensembles of InAs/InP quantum rod nanowires emitting in the telecom wavelengths[END_REF][START_REF] Anufriev | Piezoelectric effect in InAs/InP quantum rod nanowires grown on silicon substrate[END_REF]. The NWs are grown with a V/III beam equivalent pressure (BEP) ratio of 19. The growth begins with 2 min of InP at 460 °C. A growth interruption of 5 s is realized by closing the indium shutter. During this interruption, the phosphorus flux is switched to arsenic flux. Then, the growth of the InAs insertion is done by opening the indium shutter for a growth time of 20 s at 500 °C. After a 5 s growth interruption for the arsenic/ phosphorus flux switching, the final InP axial growth is done by opening the indium shutter for a growth duration of 2 min at 460 °C. The structure is completed with 10 min of InP radial growth at 340 °C. For the photoluminescence (PL) measurements, the samples are mounted in a closed-cycle temperature-controlled He cryostat, where the temperature can be varied from 14 to 300 K, and excited by a continuous-wave off-resonant excitation. The spectra are collected by a liquid nitrogen-cooled InGaAs detector with a cutoff wavelength at 2.2 μm using a conventional lock-in technique.

Results and discussion

The morphology and structure of the QR-NWs are studied by transmission electron microscopy (TEM) imaging. Fig. 1(a) shows a schematic of the investigated InAs/InP QR-NWs grown on the Si substrate. The NW diameter and length are in the 80-100 nm and 1.5-2 μm range, respectively. The length and diameter of the InAs QRs are in the range of 45-135 nm and 8-12 nm, respectively. These NWs have a WZ structure with the [0001] axis (c-axis) along the NW growth direction for both InP and InAs materials (Fig. 1(e)). Due to the high V/III BEP ratio, no stacking faults or cubic segments were observed in the QR-NWs (Fig. 1(b-e)). Similar structures are studied by SEM and show the typical four <111> directions emerging from the (001) Si [START_REF] Gao | Self-catalyzed epitaxial growth of vertical indium phosphide nanowires on silicon[END_REF][START_REF] Roest | Position-controlled epitaxial III-V nanowires on silicon[END_REF]: four directions are forming a 35°angle with the Si surface and a 90°angle from each other, and this result was also observed by our group [START_REF] Khmissi | InAs/InP nanowires grown by catalyst assisted molecular beam epitaxy on silicon substrates[END_REF]. In order to improve our understanding of the optical properties of such InAs/InP QR-NWs, low temperature PL spectroscopy experiments under different excitation powers were performed (Fig. 2).

At low excitation power, the PL spectrum shows a single peak around 0.87 eV (peak A) attributed to the ground state emission of the InAs QRs. The full width at half maximum (FWHM) is about 60 meV for this peak. The low temperature inhomogeneous broadening of the spectra is associated with the InAs QR size dispersion, which is consistent with the TEM results. When the excitation power is increased, a second peak (peak B) appears at higher energy (~0,95 eV). According to our TEM studies, the QR diameter D is 4-14 times shorter than the QR length. The direct consequence of this specific geometry is that the quantum confinement of the first electron and hole states (called ground states) are strongly related to the QR diameter. In QRs, the energy spacing between the ground states and the first excited states is small when L » D. This is confirmed by the work of different groups which have already investigated the electron and hole states in such QRs (which are also called quantum posts or columnar QDs) [START_REF] Krenner | Quantum posts with tailored structural, electronic and optical properties for optoelectronic and quantum electronic device applications[END_REF][START_REF] Nedzinskas | Electron energy spectrum in cylindrical quantum dots and rods: approximation of separation of variables[END_REF][START_REF] Andrzejewski | Eight-band k.p calculations of the composition contrast effect on the linear polarization properties of columnar quantum dots[END_REF]. Moreover, a 70 meV energy spacing has been predicted for a cylindrical InAs/InP QR with a diameter of 9.6 nm and a height of 18 nm [START_REF] Zieliński | Fine structure of light-hole excitons in nanowire quantum dots[END_REF]. As a consequence, the 80 meV energy spacing between the peak A and peak B cannot be related to the energy spacing between the ground and first excited states in our D ˜10 nm and L > 45 nm QRs.and another explanation has to be found.

In order to get additional information on the origin of this peak, the WZ InAs/InP QR-NWs have been numerically investigated by finite element method (FEM) by solving the single band effective mass Schrodinger equation in cylindrical coordinates taking into account the strain effects [START_REF] Souaf | Investigation of the InAs/GaAs quantum dots' size: dependence on the strain reducing layer's position[END_REF][START_REF] Nedzinskas | Electron energy spectrum in cylindrical quantum dots and rods: approximation of separation of variables[END_REF]. The nnth transition energy can be obtained as follow:

(1) where , are the electron and hole confined energies, respectively.

is the A or B strained InAs band gap energy. The strain dependence of is given by Ref. [START_REF] Hajlaoui | First-principles density functional theory study of strained wurtzite InP and InAs[END_REF][START_REF] Peng | Shock-induced band-gap shift in GaN: anisotropy of the deformation potentials[END_REF][START_REF] Chuang | K⋅p method for strained wurtzite semiconductors[END_REF]:

(

where is the A unstrained band gap energy and and are the deformation potentials of WZ InAs. and are the strain components parallel and perpendicular to the c-axis, respectively. If the QR is approximated by a cylinder, the strain components can be described by the Eshelby model [START_REF] Ferrand | Strain in crystalline core-shell nanowires[END_REF][START_REF] Eshelby | The determination of the elastic field of an ellipsoidal inclusion, and related problems[END_REF][START_REF] Eshelby | The elastic field outside an ellipsoidal inclusion[END_REF]: [START_REF] Gibson | Journal of Luminescence xxx (xxxx) xxx-xxx broadband high-speed single-photon detection[END_REF] where is the lattice mismatch perpendicular to the NW axis defined by were and are the lattice constant of InP and InAs, respectively. The quantity is defined by:

, where λ and μ are the Lame coefficients. For an isotropic crystal, and . represents the shape factor where D is the diameter and L the length of the QR. is given by the following analytical formula [START_REF] Artioli | Magnetic polaron in (Cd,Mn)Te quantum dot inserted in ZnTe nanowire[END_REF]:

(4) Fig. 3(a) shows the values of and as a function of the QR aspect ratio (L/D). Finally, the energy difference between A and B band energies, is deduced from the following equation [START_REF] Peng | Shock-induced band-gap shift in GaN: anisotropy of the deformation potentials[END_REF][START_REF] Chuang | K⋅p method for strained wurtzite semiconductors[END_REF]: [START_REF] Bai | InAs/InP core/shell nanowire gas sensor: effects of InP shell on sensitivity and long-term stability[END_REF] where , are the crystal field and spin orbit splittings. The value of the deformation potentials, crystal field, spin orbit splitting and unstrained band gap energies used in the calculations are listed in Table 1. All these values are theoretical values with the exception of the WZ InAs band gap being experimentally measured in Refs. [START_REF] Rota | Bandgap energy of wurtzite InAs nanowires[END_REF]. The parameters used in the energy level calculations are given in Table 2, where are the effective masses for the electron, A-band hole and B-band hole in units (free electron mass), respectively. According to the selection rule in NWs with WZ structure, the emission from recombination of electrons in the conduction band with holes from A-band is expected to be polarized perpendicularly to the NW axis, while for the recombination with holes from B-band, both perpendicular and horizontal polarizations are allowed [START_REF] Gadret | Valence-band splitting energies in wurtzite InP nanowires: photoluminescence spectroscopy and ab initio calculations[END_REF][START_REF] Anufriev | Polarization properties of single and ensembles of InAs/InP quantum rod nanowires emitting in the telecom wavelengths[END_REF]. In that case, the effective masses perpendicular to the NW axis are used in our calculations.

U N C O R R E C T E D P R O O F
The ground state energies of A and B transitions as a function of the InAs QR diameter and length are represented in Fig. 4(b).

The transition energy is found to exhibit a strong variation up to a QR length around 15 nm. However, no major contribution is observed for longer QRs. When the QR length increases from 2 nm to 60 nm, the energy spacing between the A and B transition energies decreases from ≈114 meV to ≈23 meV. Whereas a transition from heavy hole-like to light hole-like ground state is predicted in ZB InAs QRs [START_REF] Zieliński | Fine structure of light-hole excitons in nanowire quantum dots[END_REF], in WZ structures the ground state emission is always related to A-band holes. By comparing the calculated and the experimental peak energies, a QR diameter of about 10-14 nm and an average QR length of more than 15 nm can be deduced. In that case, InP NWs with larger QR diameter greater than 10 nm (Fig. 1(b)-1(c)) dominate the PL spectra. As far as the theoretical results are concerned, several phenomena could explain the discrepancy: the piezoelectric field [START_REF] Anufriev | Piezoelectric effect in InAs/InP quantum rod nanowires grown on silicon substrate[END_REF], the strain dependence of the effective masses [START_REF] Shikanai | Biaxial strain dependence of exciton resonance energies in wurtzite GaN[END_REF] and the need to take into account the exciton binding energy. Furthermore, a discrepancy is observed in the literature on the theoretical values of some fundamental parameters of WZ InAs [START_REF] Hajlaoui | First-principles density functional theory study of strained wurtzite InP and InAs[END_REF][START_REF] Rota | Bandgap energy of wurtzite InAs nanowires[END_REF][START_REF] De | Predicted band structures of III-V semiconductors in the wurtzite phase[END_REF][START_REF] Bechstedt | Structure, energetics, and electronic states of III-V compound polytypes[END_REF].

To ensure efficient optoelectronic device operating at room temperature, a clear understanding of the effect of temperature on the optical properties of such structure is highly required. Fig. 5 shows the PL spectra of the sample from 14 K to 300 K using low laser excitation power.

The PL emission peak is shifted toward higher wavelength, up to 1.55 μm at room temperature with a PL intensity keeping as much as 30% of its value at 14 K. This figure is close to the best values (50-67%) reported for ZB InAs QDs or quantum dashes grown in an InP matrix [START_REF] Lévesque | Temperature dependence of the photoluminescence spectra from InAs(P)/InP multilayers containing thick quantum dots: dot-size-dependent carrier dynamics[END_REF][START_REF] Hostein | Time-resolved characterization of InAsPInP quantum dots emitting in the C -band telecommunication window[END_REF]. As the temperature increases, the QR-NW emission peak due to the thermal escape of carriers from A-band to B-band. A multi-Gaussian fitting is performed on each PL spectrum to take into account the fact that two peaks are observed on the spectrum at room temperature with an average energy spacing of 70 ± 5 meV which is close to the average energy splitting observed between the A-and B-bands at low temperature.

The temperature dependence of the A exciton integrated PL intensity is shown in Fig. 6(a). A fit of the experimental data has been performed using the following multi-channel Arrhenius equation [START_REF] Bando | Narrowing of exciton linewidth of a quantum dot with increasing temperature[END_REF]: [START_REF] Hill | Interplay of surface recombination and diode geometry for the performance of axial p-i-n nanowire solar cells[END_REF] where i represents the number of involved channels for carrier thermal escape, is the Boltzmann constant, is the thermal activation energy, is the PL intensity at 0 K and is a constant. We found that the temperature dependence of the integrated PL intensity can be well fitted by considering two escape channels (i = 2). The best fit gives the values of E 1 and E 2 around 10 (±2 meV) and 125 (±10 meV) meV, respectively. These values are comparable with the findings of Wu et al. [START_REF] Wu | Defect-free self-catalyzed GaAs/GaAsP nanowire quantum dots grown on silicon substrate[END_REF] were 14.3 and 73.4 meV activation energies were obtained for GaAs/GaAsP NW quantum dots and associated with the GaAs surface states, which limits the optical properties of the NWs, and to thermal excitation of photocarriers from the GaAs QDs into the GaAsP barriers, respectively. Alonso-Álvarez et al. [START_REF] Alonso-Álvarez | Strain balanced quantum posts[END_REF] also reported 11 and 170 meV activation energies for InAs/GaAsP quantum posts (QPs). The lower activation energy is associated to carrier recombination through defects in the quantum structures or their interfaces while the higher activation energy is compatible with the unipolar escape of carriers from the nanostructures to the matrix [START_REF] Alonso-Álvarez | Strain balanced quantum posts[END_REF]. If we assume an isotropic dielectric constant, the WZ InAs exciton binding energy can be estimated using the following equation: [START_REF] Wagner | Vapor-liquid-solid mechanism of single crystal growth[END_REF] where ε(0) is the dielectric constant (15.15 for ZB InAs [START_REF] Tang | Binding energy of the free exciton in indium arsenide[END_REF]), m e and m A are the electron and A-band hole effective masses that takes in account the existing anisotropy in the WZ InAs effective masses. Using the parameters given in Table 2, we thus obtain E b = 2 meV. This value does not differ from that of the free exciton in bulk ZB InAs (1 meV) [START_REF] Tang | Binding energy of the free exciton in indium arsenide[END_REF]. Moreover, the exciton binding energy is in the 5-25 meV range for SK InAs/InP QDs [START_REF] Pettersson | Electrical and optical properties of self-assembled InAs quantum dots in InP studied by space-charge spectroscopy and photoluminescence[END_REF][START_REF] Cornet | Impact of the capping layers on lateral confinement in InAsInP quantum dots for 1.55 μm laser applications studied by magnetophotoluminescence[END_REF]. In that case, the smaller activation energy E 1 = 10 meV can be associated to the exciton binding energy in WZ InAs. However, the high temperature (160-300 K temperature range) activation energy E 2 is about five times less than the difference (≈620 meV) between the energy of the WZ InP barrier band gap (1.49 eV) and the energy of the PL emission of the A transition (0.87 eV). In the other hand, this value is close in magnitude to the energy spacing between the A-and B-bands (80 meV). In that case, E 2 can be associated to the thermally populated high-energy states in the B-band. This thermal population is confirmed by the relative increase of B transition PL intensity along with the temperature as observed for WZ InP NWs [START_REF] Hadj Alouane | Excitonic properties of wurtzite InP nanowires grown on silicon substrate[END_REF]. Fig. 6(b) shows the FWHM of the A peak as a function of temperature. The FWHM almost remains constant for temperature below 100 K. With further increasing temperature to 160 K, the FWHM was slightly increased by about 5-6 meV and then remains constant at higher temperature range. This behavior has been observed for SK InAs/GaAs QDs and attributed to the thermal population of excited hole states [START_REF] Heitz | Temperature dependent optical properties of self-organized InAs/GaAs quantum dots[END_REF]. As far as our sample is concerned, this behaviour might be attributed to the thermal population of the B-band. This is in a good agreement with studies on the temperature dependence of the A transition integrated PL intensity.

U N C O R R E C T E D P R O O F
The temperature dependence of the emission energy of the A peak is plotted in Fig. 6(c). Generally, this behaviour is discussed in terms of the empirical equation of Varshni [START_REF] Varshni | Temperature dependence of the energy gap in semiconductors[END_REF] and Bose-Einstein formula [START_REF] Singh | Temperature dependence of the lowest excitonic transition for an InAs ultrathin quantum well[END_REF] given by: (8) [START_REF] Tchernycheva | Growth and characterization of InP nanowires with InAsP insertions[END_REF] where α and β parameters are the Varshni coefficients, aBE represents the strength of the electron-phonon interaction and is the average phonon temperature. Solid and dashed lines in Fig. 6(c) are the fitting curves of the experimental data based on equations ( 8) and ( 9), respectively. The best fitting parameters for the two models are summarized in Table 3 along with parameters from the literature. These parameters are closed to the values reported in the literature for ZB InAs which indicates that ZB and WZ InAs have similar thermal parameters as observed for InP [START_REF] Hadj Alouane | Excitonic properties of wurtzite InP nanowires grown on silicon substrate[END_REF][START_REF] Zilli | Temperature dependence of interband transitions in wurtzite InP nanowires[END_REF] and GaAs [START_REF] Vainorius | Temperature dependent electronic band structure of wurtzite GaAs nanowires[END_REF] (even if there is some discrepancy for this semiconductor materiel [START_REF] Ahtapodov | A story told by a single nanowire: optical properties of wurtzite GaAs[END_REF]). The low temperature data points (<100 K) are systematically below the Varshni curves by about 5 meV. This value is close in magnitude to the smaller activation energy E 1 calculated from the fitting of the temperature dependence of the integrated PL intensity by the Arrhenius equation and attributed to the formation of excitons at low temperatures. When the temperature is raised further, the energy shift follows the Varshni law. This constitutes an additional proof of the good quality of our InAs/InP QR-NWs grown on silicon.

Conclusions

In summary, we have grown high quality InAs/InP QR-NWs on Si(001) substrates by VLS assisted MBE. Optical studies performed at cryogenic temperature reveal an InAs QR emission in the telecom band with two peaks. FEM calculations taking into account strain effects in the InAs QR were developed to explain these experimental results. Temperature-dependent optical properties revealed a high stability of the PL intensity in the 14-300 K temperature range and showed an emission at the 1.55 μm wavelength at room temperature. These results confirm a great promise for InAs-InP QR-NWs for telecom-band na noemitters monolithically integrated on silicon substrates for photonic chips. This study 270 [START_REF] Chouaib | Photoreflectance spectroscopy of self-organized InAs/InP(0 0 1) quantum sticks emitting at 1.55 μm[END_REF] 99 [START_REF] Chouaib | Photoreflectance spectroscopy of self-organized InAs/InP(0 0 1) quantum sticks emitting at 1.55 μm[END_REF] 271 [START_REF] Paul | Empirical expressions for the alloy composition and temperature dependence of the band gap and intrinsic carrier density in GaxIn 1-xAs[END_REF] 4.19 [START_REF] Paul | Empirical expressions for the alloy composition and temperature dependence of the band gap and intrinsic carrier density in GaxIn 1-xAs[END_REF] 247 [START_REF] Singh | Temperature dependence of the lowest excitonic transition for an InAs ultrathin quantum well[END_REF] 56 [START_REF] Singh | Temperature dependence of the lowest excitonic transition for an InAs ultrathin quantum well[END_REF] 93 [START_REF] Vurgaftman | Band parameters for III-V compound semiconductors and their alloys[END_REF] 2.76 [START_REF] Vurgaftman | Band parameters for III-V compound semiconductors and their alloys[END_REF] U

N C O R R E C T E D P R O O F

Fig. 3 (

 3 b) shows the QR aspect ratio dependence of the A and B band gap energies. For infinite QR , the energy difference between A and B band gap energies remains constant. In the other hand, for flat QR , the A and B band gap energies are strongly affected by the strain effect. In our case, the calculations are based on WZ InP NWs with infinite length and diameter. Fig. 4(a) shows the schematic structure of the band alignment of the WZ InAs QR inserted in a WZ InP NW used in our calculations.

Fig. 1 .

 1 Fig. 1. (a) Schematic of the investigated sample. (b-d) TEM images showing InP NWs with different dimensions of InAs QRs. (e) A close-up HAADF-TEM image of the InAs/InP interface.

Fig. 2 .

 2 Fig. 2. Normalized PL spectra at 14 K of InAs/InP QR-NWs as a function of the excitation power. The solid lines are the experimental data and the dotted lines are the Gaussian fit of the PL spectra.

Fig. 3 .

 3 Fig. 3. (a) Strain components and (b) Energies of the strained A and B band gap energies, as a function of the QR aspect ratio (L/D).

Fig. 4 .

 4 Fig. 4. (a) Schematic representation of the band alignment for WZ InAs/WZ InP used in our calculation. (b) Simulated A and B transition energies of WZ InAs/InP QR-NWs as a function of the QR diameter (D) and length (L) (from 2 nm to 60 nm). Horizontal dashed lines represent the energies of A and B peaks observed by PL.

Fig. 5 .

 5 Fig. 5. Temperature induced changes in the PL spectrum for InAs/InP QR-NWs on Si(001) using low laser excitation power. The solid lines are the experimental data and the dotted lines are the multi-Gaussian fitting of the room temperature PL spectra. The vertical dashed line indicates the 1.55 μm emission wavelength.

Fig. 6 . 3

 63 Fig. 6. Temperature dependence of the: (a) integrated PL intensity, (b) FWHM and (c) PL peak energy of A exciton (green circles). The solid curves give the fitting result.

Table 1 The

 1

	deformation potentials	and , the crystal field splitting	, spin orbit splitting	and the unstrained band gap energy	(eV) used in calculations.
				(eV)	(eV)	(eV)
	0.379 [36]	0.195 [36]	0.477 [35]	-3.17 [29]	5.85 [29]	-8.33 [29]	-2.69 [29]	InAs (WZ)

Table 2

 2 WZ InAs and WZ InP effective masses and lattice parameters used for the calculations. and are the effective masses perpendicular and parallel to the c-axis, respectively.

		c (Å)	a (Å)
	InP (WZ) InAs (WZ)	6.77 [29] 7 [29]	4.15 [29] 4.283 [29]	0.088 [36] 0.042 [36]	0.158 [36] 0.084 [36]	1.273 [36] 1.7 [36]	0.169 [36] 0.113 [28]	0.839 [36] 0.101 [36]

Acknowledgments

This work has been supported by the French Agence Nationale pour la Recherche (ANR) "INSCOOP" project. The authors gratefully thank the Nanolyon plateform for access to equipments and J. B. Goure, C. Botella, P. Regreny, P. Cremillieu, and R. Mazurczyk for technical assistance. The authors acknowledge the Deanship of Scientific Research at King Faisal University for financial support under the Research Group Support Track (Grant No. 17122003).

Author statement

Mohamed Helmi Hadj Alouane, Nicolas Chauvin and Catherine Bru-Chevallier carried out the PL measurements on QR-NWs, data analyses, strain calculations and wrote the major part of the paper. Olfa Nasr, Mohamed Hadj Alouane, Mohamad Mahmoud Ahmad and Bouraoui Ilahi contributed in data analyses and FEM calculations. Hammadi Khmissi and Michel Gendry conducted the MBE growth. Gilles Patriarche performed TEM measurements. All authors reviewed the manuscript.

Declaration of competing interest

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


I S Loginova
email: loginova@misis.ru

M V Sazerat

P A Loginov

A V Pozdniakov

N A Popov

A N Solonin

Evaluation of Microstructure and Hardness of Novel Al-Fe-Ni Alloys with High Thermal Stability for Laser Additive Manufacturing

The microstructure and phase composition of cast and laser-melted Al-Fe-Ni alloys were investigated. Two main phases-Al 3 (Ni,Fe) and Al 9 FeNi-were formed in the as-cast state. A fine microstructure without porosity or solidification cracks was observed in the Al-Fe-Ni alloys after laser treatment. The hardness of the laser-melted alloys was 2.5-3 times higher than the hardness of the as-cast alloys owing to the formation of an aluminum-based solid solution and fine eutectic particles. The formation of the primary Al 9 FeNi phase was suppressed as a result of the high cooling rate. Annealing these alloys at temperatures less than 300°C demonstrated the high thermal stability of the microstructure while maintaining the hardness. The Al-Fe-Ni alloys investigated in this study are promising heat-resistant materials for additive manufacturing because of their fine, stable structure, and the low interdiffusion coefficients of Fe and Ni.

INTRODUCTION

Additive manufacturing (AM) technologies offer the capability to produce parts with improved properties. Selective laser melting (SLM) and direct metal deposition (DMD) are the most common technologies used to manufacture parts from metal powders. In both technologies, parts are formed layer by layer from the powder upon laser beam impact. [1][2][3] Aluminum (Al) alloys, having low density and desirable mechanical and technological properties, are good candidates for AM.

The list of Al alloys suitable for AM is increasing every year and currently includes powders in the following forms: Al-Si (AlSi12, 4 AlSi10Mg, 5 AlSi7Mg 6 ), Al-Mg-Sc (Al-6.2Mg-0.36Sc-0.09Zr, 7 Al-4.6Mg-0.66Sc-0.42Zr-0.49Mn 8 ), Al-Zn-Mg (AA7075, 9 AA7050 10 ), and Al-Cu-Mg (Al-3.5Cu-1.5Mg-1Si 11 ). However, only AlSi10Mg alloy and Al-4.6Mg-0.66Sc-0.42Zr-0.49Mn alloy (Scalmalloy) powders are currently being used in commercial applications. The latter material is conducive to the laser melting process because of its narrow effective solidification range and the presence of a sufficiently large amount of aluminum-silicon eutectic in the structure, as well as the high content of zirconium and scandium. The AlSi10Mg alloy has substantially lower strength than the Scalmalloy but is significantly less expensive. The Al-Cu and Al-Cu-Mg alloys exhibit relatively high strength at both room and elevated temperatures. [START_REF]Properties and Selection: Nonferrous Alloys and Special-Purpose Materials[END_REF] However, the drawback of using these alloys in AM lies in their poor casting properties and weldability. [START_REF] Lopez-Botello | [END_REF][14][15][16] Aluminum alloys with transition metals are prospective heat-resistant material options for applications at 300-350°C. 2,[17][18][19][20] The development of novel material compositions with strong mechanical and technological properties, especially at elevated temperatures, is currently a focus point for AM. Alloying with transition metals such as Mn, Cr, Fe, Ni, and Ce has been shown to increase hightemperature strength and improve the casting properties of Al alloys. 2,16,18,19 A negative outcome when alloying with transition metals is the formation of coarse primary intermetallic particles. However, the ultra-high cooling rate of 10 3 -10 5 K/s helps to inhibit the nucleation and recrystallisation of primary intermetallic particles. [21][22][23] Another advantage of the high cooling rate is the potential formation of a supersaturated solid solution (SSSS). Stability of the SSSS is one of the most important factors for heat-resistant alloys. The stability of both the SSSS and intermetallic phases depends on the rate of diffusion of the alloying elements in the Al solid solution. Good candidates for Al alloying are Fe and Ni because of their low diffusion coefficients (7 9 10 À13 and 2 9 10 À1224 in Al at 500°C, respectively), which are less than one-third of the diffusion coefficients of the main alloying elements (Mg, Cu, Zn, Si).

Alloying of Al with Fe and Ni has attracted great interest in this system in applications of AM, in particular for selective laser melting. Slow-cooled Al-Fe-Ni alloys are characterized by coarse primary brittle intermetallic phases. 25 Such structural features may impair the mechanical properties of these materials, confining their application. Hence, a high cooling rate, that can be realized in AM, can solve this problem. Alloying with nickel and iron improves the high-temperature mechanical properties 26 of the Al 3 Fe, Al 3 Ni and Al 9 FeNi intermetallic phases. 27 A high Ni content is beneficial in cases of selective laser melting because the formation of Al 9 FeNi primary crystals can be suppressed.

The development of new alloys for AM technologies is associated with the complexity of the manufacture of the powders, the optimization of the production technology, and the high cost of the equipment. However, the development of powders is a classical material science task, in which it is important to understand the structure formation processes after laser treatment. The impact is characterized by the following features: high local temperature leading to the melt of a small volume of metal, and subsequent fast cooling under conditions of directional heat removal. 7 These features determine the basic principles of the formation of the structure and properties of products produced by selective laser melting and direct laser deposition. 20 These features can be simulated on a simple laser machine using bulk samples 18 that allow the analysis of the basic principles of structure formation of the alloys without the application of powder materials.

In the present work,the laser melting of Al-Fe-Ni bulk metal was used as a method of simulation of SLM and DLD melting and solidification conditions to analyze the formation of the microstructure of laser-melted Al-Fe-Ni alloys during solidification and heat treatment.

MATERIALS AND METHODS

Novel Al-2.5Fe-5.5Ni, Al-2.5Fe-7.5Ni, and Al-2.5Fe-9.5Ni alloys were prepared by melting a mixture of pure Al (99.99%), Al-10Fe master alloy, and Al-20Ni master alloy in a Nabertherm S3 electric furnace at 780°C. In order to check the chemical composition of the alloys, we carried out energy-dispersive x-ray spectroscopy (EDS) analysis of the large areas of polished samples: 1 9 1 mm. The real composition was practically identical: Al-2,4%Fe-5,7%Ni, Al-2,5%Fe-7.6%Ni, and Al-2,9%Fe-9,3%Ni. These concentrations of alloying components were chosen after the analysis of the ternary phase diagram of the Al-Fe-Ni system: all the alloys were located in the phase region: (Al), Al 3 (Fe,Ni), and Al 9 FeNi. As the Ni content increases, the fraction of the heat-resistant Ni-containing phase Al 3 (Fe,Ni) should also increase, leading to increases in the strength and heat resistance of the alloys. Although the Al 9 FeNi phase forms as coarse crystals under equilibrium conditions, in fast crystallization during laser melting, the Al 9 FeNi formation can be suppressed leading to the formation of SSSS with high mechanical properties.

Ingots with the size of 40 9 20 9 120 mm 3 were cast into a water-cooled copper mold. The cooling rate of the as-cast ingot was approximately 10 K/s. The cooled ingots were cut into 2-mm-thick plates for laser melting.

The liquidus and solidus temperatures were determined by differential scanning calorimetry (DSC) on a Setaram Labsys calorimeter in an argon atmosphere. The time dependence of the temperature difference between the cell with the sample and the reference cell was measured experimentally using an S-type (platinum:platinum-rhodium) thermocouple. The measurement error was ± 1°C. The experiments were carried out in the 20-800°C temperature range at a constant heating rate (1°C/ s).

Laser melting of single tracks was carried out on the 2-mm-thick plates using an MUL-1-M-200 machine equipped with an Nd-YAG fiber laser (1064 nm wavelength) under the following parameters: laser power 150 W, scan speed 1 mm/s, overlap distance 0.2 mm, and an argon atmosphere. These parameters provided an average volumetric energy density of approximately 50 J/mm 3 . The alloys were heat-treated in a Nabertherm N30/65A electric furnace.

The volume of the eutectic phases in the as-cast state was determined by using Axis Visual software. The size of the dendritic cells was measured using the linear intercept method according to the Russian standard 21073.0-75.

To determine the phase composition and to calculate the lattice parameter in the specimens that were laser-melted and annealed at 500°C for 1 h, xray powder diffraction (XRD) analysis was performed using a Brucker Advance D8 x-ray diffractometer. For the XRD studies, we prepared special square specimens of 1 mm thickness. The specimens were laser-melted through their whole surface and thickness. The XRD results of the laser-melted and laser-melted/annealed specimens were provided to confirm that the Al 9 FeNi phase was suppressed.

Annealing at 500°C for 1 h was chosen to complete the solid solution decomposition, which was observed by means of transmission electron microscopy (TEM; JEOL 2100) and energy-dispersive xray spectroscopy (EDX; XMAX-80). The microstructure, chemical composition, and phase composition of the alloys were analyzed by scanning electron microscopy (SEM; TESCAN VEGA 3LMH) in conjunction with EDS. The structures of the foils were investigated by TEM at 200 kV. Specimens with a diameter of 3 mm were prepared using the A2 electrolyte on Struers Tenupol-5 equipment.

To analyze the heat resistance of the investigated alloys in laser-melted condition the next regimes were used: annealing at 100°C, 200°C, 250°C, 300°C, 350°C, 400°C, and 500°C for 1 h, at 250°C for 1, 2, 3, 4, 8, 16, and 20 h, and at 300°C for 1, 2, 3, 4, 8, 16, and 20 h.

Vickers hardness (HV) measurements were performed on a Wolpert 402MVD hardness tester.

RESULTS AND DISCUSSION

The microstructures of the alloys in the as-cast state are presented in Fig. 1. The dendrites of the Al-based solid solution and of the primary crystals of Al 9 FeNi phase are the main microstructural elements in the as-cast state. The two phases that form the secondary eutectic were identified as Al 3 (Ni,Fe) and Al 9 FeNi (30 9 400 lm 2 grains in Fig. 1d). As the Ni concentration increased from 5.5% to 9.5%, the volume percentage of the eutectic increased from 24% to 37% (Fig. 1a-c).

For the alloys used in AM, the solidification range is extremely important. [28][29][30] Solidification cracking is closely related to the range of the solidification temperature and the content of liquid phases. In the current work, the comparison between the solidification intervals of Al-Fe-Ni and AlSi10Mg (the most popular alloy in SLM 20 ) alloys were carried out (Table I). Liquidus and solidus temperatures and the solidification range of the Al-Fe-Ni alloys were detected by DSC and are presented in Table I.

The investigated alloys have a narrow solidification range that can provide good casting properties and weldability. Also, the Al-Fe-Ni alloys have much higher solidus temperatures than AlSi10Mg (Table I) which provide high heat resistance.

A typical SEM image of the microstructure of the laser-melted single tracks tested is shown in Fig. 2. A structure without porosity and solidification cracks was formed in the Al-Fe-Ni alloys being investigated. The formation of the primary Al 9 FeNi phase was suppressed owing to the high cooling rate. Figure 3 provides TEM images of the microstructure of the laser-melted single tracks. An Al-based solid solution and fine eutectic origin particles were present in the microstructure. EDX analyses of the TEM images confirms the formation of Al 3 (Ni,Fe) (Fig. 3d). The mean sizes of the dendritic cells of the primary Al were 165 nm, 80 nm, and 27 nm in the Al-2.5Fe-5.5Ni, Al-2.5Fe-7.5Ni, and Al-2.5Fe-9.5Ni alloys, respectively. A decrease of the dendritic cell average size in the alloys may be attributed to a change in the crystallization conditions. It has been shown that, as the Ni concentration increases, the solidification range decreases (Table I). Under the same cooling conditions, provided by one laser melting mode, the degree of supercooling will be higher in the Al-2.5Fe-9.5Ni alloy compared to the Al-2.5Fe-5.5Ni alloy, leading to the appearance of more crystallization nuclei and the refinement of the microstructure. A quasi-eutectic structure was found in the Al-2.5Fe-9.5Ni alloy (Fig. 3c), which features a similar mean size of the dendritic cells and eutectic phase particles (10-70 nm). The lattice parameters of the Al-based solid solutions in the laser-melted and annealed specimens measured via XRD are shown in Table II. The lattice parameter of the Al solid solution decreases with the Ni content from 0.40487 nm in the Al-2.5Fe-5.5Ni alloy to 0.40465 nm in the Al-2.5Fe-9.5Ni alloy. The reduction of the lattice parameter in Al is associated with the dissolution of Ni in it by the substitution method. This was presumably due to the increase in the Ni content in the Al solid solution given that the atomic radius of Ni 24 is 0.019 nm smaller than that of Al. 24 After annealing, the lattice parameters approached that of pure Al (99.99%), which is 0.40491 nm. 24 This indicates the decomposition of a supersaturated Al-based solid solution followed by the formation of a secondary Ni-containing phase. The XRD analysis of the samples after annealing showed that the intensity of the peaks of the Al 3 (Fe,Ni) phase increases (see Supplementary Fig. S1) indicating an increase in the concentration of this phase in the alloys. Hence, the secondary phase has the same structure as the eutectic.

The hardness in the as-cast state for the Al-2.5Fe-5.5Ni, Al-2.5Fe-7.5Ni and Al-2.5Fe-9.5Ni alloys are 53, 58, and 65 HV, respectively. A slight tendency for the hardness to possibly increase can be associated with an increase of the primary crystal size and the volume of the eutectic phases. The structural refinement and formation of a SSSS leads to a significant increase in hardness, by more than a factor of 3, compared to the as-cast alloys (see Supplementary Fig. S2).

Annealing of the alloys at lower temperatures (below 300°C) demonstrated the high thermal stability of the microstructure and the hardness (Figs. 4,5). HV curves of the Al-Fe-Ni alloys after annealing at 100°C, 200°C, 250°C, 300°C, 400°C, and 500°C for 1 h, 250°C for 1-20 h, and 300°C for 1-20 h are presented in Fig. 4. The hardness slightly decreased after a -h annealing at temperatures below 300°C, but fell drastically at higher annealing temperatures (Fig. 4a). Annealing at 250°C and 300°C resulted in a slight decrease in HV after 1-5 h and stabilisation of HV up to 20 h. Softening at low temperatures (e.g., 300°C) could be due to the decrease in the as-cast dislocation density and partial decomposition of the SSSS. The size and morphology of the intermetallic phases after annealing at 250°C for 20 h and 300°C for 1 h experienced no significant changes (Fig. 5a,b). The sharp decrease in hardness after a 1-h annealing at 500°C can be explained by the fragmentation and growth of the intermetallic phases (Fig. 5c), and the full decomposition of the supersaturated Al solid solution (Table II). The size of the eutectic phase particles increased to 100-300 nm after annealing at 500°C for 1 h (Fig. 5c),and the lattice parameter increased up to the value of pure Al (Tables II, III).

CONCLUSIONS

The microstructure and phase composition of ascast and laser-melted Al-Fe-Ni alloys were investigated by SEM, TEM, and XRD. Al 3 (Ni,Fe) and Al 9 FeNi phases formed in the as-cast alloys. A structure without porosity or solidification cracks was formed in the Al-Fe-Ni alloys after laser treatment. The hardness of the laser-melted alloys was 2.5-3 times higher than the hardness of the ascast alloys. An Al-based solid solution and fine eutectic particles were present in the microstructure. The formation of the primary Al 9 FeNi phase was suppressed due to the high cooling rate. The mean size of the dendritic cells of the primary Al in the Al-2.5Fe-5.5Ni, Al-2.5Fe-7.5Ni, and Al-2.5Fe-9.5Ni alloys were 165 nm, 80 nm and 27 nm, respectively. The size of the eutectic particles ranged from 10 nm to 70 nm. The Ni content in the Al solid solution increased with the Ni content in the alloy because the atomic radius of the Al is smaller than that of Ni. Annealing of the investigated alloys at lower temperatures (below 300°C) demonstrated the high thermal stability of the microstructure and the hardness. Annealing at 250°C and 300°C leads to a slight decrease in HV after 1-5 h and stabilization of the HV up to 20 h. The softening could be due to the decrease in the ascast dislocation density and the partial decomposition of a SSSS. The size and morphology of the intermetallic phases after annealing at 250°C for 20 h and 300°C for 1 h changed insignificantly. The Al-Fe-Ni alloys investigated in this study are promising high-strength and heat-resistant materials for AM because of their fine, stable structure and low diffusion rates of Fe and Ni.

Fig. 1 .

 1 Fig. 1. SEM image of (a) Al-2.5Fe-5.5Ni, (b) Al-2.5Fe-7.5Ni, and (c) Al-2.5Fe-9.5Ni alloys, and (d) XRD patterns of the Al-Fe-Ni alloys in the ascast state.

Fig. 2 .

 2 Fig. 2. SEM image of a single track of Al-Fe-Ni alloys.

Fig. 3 .

 3 Fig. 3. TEM images of single tracks in (a) Al-2.5Fe-5.5Ni, (b) Al-2.5Fe-7.5Ni, and (c) Al-2.5Fe-9.5Ni alloys, and (d) Fe and Ni distribution maps of the Al-2.5Fe-7.5Ni alloy.

Fig. 4 .

 4 Fig. 4. HV curves of Al-Fe-Ni alloys after annealing at (a) 100-500°C for 1 h, (b) 250°C for 120 h, and (c) 300°C for 1-20 h. The black line is Al-2.5Fe-5.5Ni, the red line is Al-2.5Fe-7.5Ni, and the blue line is Al-2.5Fe-9.5Ni.

Table I .

 I Liquidus and solidus temperatures and the solidification range of the Al-Fe-Ni alloys (experimental data)

	Alloy	Liquidus temperature, °C	Solidus temperature, °C	Solidification range,	°C
	Al-2.5Fe-5.5Ni	641	635	6
	Al-2.5Fe-7.5Ni	641	638	3
	Al-2.5Fe-9.5Ni	641	639	2
	AlSi10Mg 31	578	557	21

Table II .

 II Lattice parameters of Al solid solutions of laser-melted and laser-melted/annealed areas

	Alloy	Condition	A ˚DA Ål-2.5Fe-5.5Ni
		Laser-melted	4.0487 0.0008
		Laser-melted/annealed 4.0496
	Al-2.5Fe-7.5Ni	Laser-melted	4.0480 0.0015
		Laser-melted/annealed 4.0495
	Al-2.5Fe-9.5Ni	Laser-melted	4.0465 0.0030
		Laser-melted/annealed 4.0495

Table III .

 III EDS results of the Al-2.5Fe-7.5Ni alloy after annealing at 500°C for 1 h Chemical composition, at.%

		Al	Ni	Fe
	EDS 1	71	23	6
	EDS 2	100	0	0

ACKNOWLEDGEMENTS

Loginova I.S. would like to thank Dr. Solonin A.N. for valuable discussions regarding the structure formation process. This project and all the experiments were funded by RFBR, Project Number 19-38-60037.

CONFLICT OF INTEREST

The authors declare that they have no conflict of interest.

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