Direct synthesis of amorphous metal–organic frameworks from nanoclusters

Domain: neuroscience
Published: 2026-04-08
Editorial outcome: major_revision

Abstract
The direct synthesis of amorphous metal–organic frameworks (aEMOFs) is an appealing yet often avoided
approach due to the perceived unpredictability of amorphous network formation. Here, we develop a
strategy for aMOF synthesis using pre-formed nanoclusters and rigid organic linkers, providing enhanced
control over disorder and defect chemistry while bypassing the traditional crystallise–amorphise approach.
By systematically comparing this approach to other direc t synthesis routes and using X-ray pair distribution
function, thermogravimetric and statistical analyses, we establish key design principles governing aMOF
formation. We demonstrate that kinetic contIrol—fast reactions under basic conditions at room
temperature—suppresses crystallisation and drives amorphous network formation. The nanocluster
approach consistently yields highly disordeEred frameworks with the shortest coherence lengths among
direct synthesis methods. Additionally, Lwe show that tuning metal composition through doping with
kinetically inert cations restricts coordination reversibility, significantly increasing defect density and
structural disorder. Structural analysis reveals that while aMOFs share motifs with crystalline polymorphs,
they cannot be directly mapped onto known structures, highlighting the importance for more nuanced
characterisation approaches. By developing a systematic approach to aMOF design, this work provides a
foundation for tailoring strucRtural disorder, expanding their potential for catalysis, adsorption, and transport
applications.

Introduction
Crystal engineering has played a pivotal role in the rapid development of metal–organic frameworks
(MOFs), leading to a diverse range of materials with applications in gas storage, catalysis, and drug
delivery.1 Crystalline frameworks have been the primary focus due to their well-defined structure–property
relationships, which enable rational material design.2–4 In contrast, amorphous MOFs (aMOFs) are gaining
attention for their enhanced processability, catalytic properties, and greater mechanical integrity when
formed as monoliths.5–11 However, their discovery and structural characterisation remain a significant
challenge.9

Most aMOFs are produced through a crystallise–amorphise approach, in which a crystalline MOF is first
synthesised and then amorphised by heat or pressure.6,9 Depending on the system and conditions, such
treatments can also yield glassy MOFs that exhibit a glass transition, representing one subset within the
broader family of non-crystalline MOF materials.10 These methods pose a fundamental bottleneck in aMOF
discovery, limiting their structural diversity by restricting them to derivatives of crystalline frameworks. As a
result, aMOFs are often perceived merely as disordered analogues of crystalline MOFs.6 To fully realise
their potential as functional materials, synthetic strategies that bypass the need for a crystalline MOF are
essential.

Direct synthesis offers a promising alternative, as it avoids the need for a crystalline MOF precursor.12
However, it is overlooked due to challenges in controlling the formation of amorphous networks. The kinetic
nature of aMOFs and their relatively flat energy landscapes can lead to unpredictable structures, making
rational design of direct synthesis difficult. Furthermore, polymorphism—a well-known challenge in MOFs—
is even more complex in aMOFs, where distinguishing between different polyamorphic phases is
notoriously difficult yet can have a significant impact on material properties. For example, computational
studies on the Fe‑BTC (BTC: 1,3,5-benzentricarboxylate) aMOFs have shown that variations in local order
and defect concentrations independently influence properties such as pore size and surface area.13,14 While
a small number of studies have demonstrated that direct synthesis of aMOFs is possible by kinetically
suppressing crystallisation, these examples remain comparatively rare and system-specific, underscoring
that generalisable design rules for the rational, reproducible synthesis of aMOFs are still largely lacking.12,15

To address this, we focus on the related Fe‑BDC family, a system known for its extensive crystalline
polymorphism.16 By reacting an Fe salt with 1,4‑benzenedicarboxylic acid (H BDC), multiple crystalline
2
polymorphs can be formed, including MIL‑101 (MIL; Material Institute Lavoisier), MIL‑88, MIL‑68, MIL‑53,
MIL‑235, and Fe(BDC)(DMF).16–21 These phases exhibit distinct local structures, topologies, and
responsive behaviour. In 2022, an amorphous Fe‑BDC material with promising catalytic properties was
reported and proposed as the amorphous analogue of the crystalline SMIL‑101 framework.17,22 However,
given the extensive polymorphism within Fe‑BDC materials and the limited structural characterisation
provided, it remains unclear whether this aMOF can be definitively assigned as the amorphous counterpart
to MIL‑101. This raises the broader question: what defines theE amorphous counterpart to a crystalline
material? It is a question that has long been posed across fRamilies of important materials, such as oxide
glasses, metallic glasses, and amorphous semiconductors.23–25 This issue becomes especially significant
when amorphous materials are synthesised directly, and therefore, no obvious parent crystalline material
is present. Addressing this question requires a systema tic approach to direct aMOF synthesis.

One strategy that successfully expanded the crysItalline MOF field—particularly for complex, hierarchically
structured frameworks—was the use of pre-syn thesised secondary building units (SBUs).26–29 By directly
reacting discrete SBUs or nanoclusters withE organic linkers, greater control over framework connectivity
and topology can be achieved. This strLategy has already enabled the synthesis of crystalline Fe-BDC
polymorphs such as MIL-88.18 While dCirect synthesis of aMOFs from nanoclusters has been demonstrated
using highly flexible linkers to kinetically suppress crystallisation, the combination of nanoclusters with rigid
organic linkers to generate aMOFs remains essentially unexplored.30

In this work, we challengeR the crystallise–amorphise paradigm by developing a nanocluster-based
approach for the direct sAynthesis of aMOFs, specifically non-crystalline MOFs without a glass transition.
We synthesise a series of M-BDC materials (M = Fe, Fe₂Cr, FeCr₂, and Cr) to investigate how synthesis
conditions and chemical composition influence aMOF structure. Through a systematic study comparing
nanocluster-based synthesis to one-pot and stepwise approaches using X-ray pair distribution function,
thermogravimetric, and statistical analyses, we establish fundamental design principles for aMOF
formation. Our goal is simple yet unconventional: To produce aMOFs with (i) well‑defined local structure,
(ii) minimal long‑range order, and (iii) tuneable defect concentrations. By targeting highly disordered
frameworks, we seek to maximise sample isotropy, a property that has the potential to enhance mass
transport, while controlling defect concentrations can be leveraged for improved catalytic performance.15,31
This study expands the structural diversity of aMOFs, moving beyond current limitations and providing a
rational framework for their design.

Results
We begin by providing an overview of our approach and outlining the key sections of this study. Our primary
focus is to develop the nanocluster (NC) approach for aMOF synthesis and compare it with one‑pot (OP)
and stepwise (SW) methods [Fig. 1a]. We thus compare three different synthesis methods: In the NC
approach, a pre‑synthesised nanocluster is used as a reagent that is reacted directly with an organic linker
in solution. In the OP approach, the metal salt is directly combined with the organic linker in solution. The
SW approach is an intermediate between the NC and OP methods: the metal salt is first reacted with
acetate ions to form the nanocluster in solution, which is then reacted with the organic linker.

For this study, we selected tRhe M₃O trimer as our nanocluster due to its role as an archetypal SBU in the
MOF field [Fig. 1b]. TheA M₃O trimer is a fundamental building block of many canonical MOFs, such as
MIL‑100, MIL‑101, and MIL‑53.16,17,26,32 Importantly, it can be readily synthesised as a discrete nanocluster
([M O(CH CO ) (H O) ]NO ) using acetate groups as capping ligands, making it well‑suited for the NC
3 3 2 6 2 3 3
approach.26 The M₃O trimer has been extensively studied, and its ability to incorporate multiple metal types
enables us to investigate both the chemical composition and synthetic method in the direct synthesis of
aMOFs.33–38 We specifically selected Fe and Cr for their similar ionic radii (0.645 and 0.615 Å for Fe3+ and
Cr3+, respectively), enabling the complete compositional series to be synthesised, and their different
chemical lability.39

This study is ordered as follows: (i) nanocluster synthesis, including formation kinetics and atomic structure,
(ii), optimisation of the NC approach for aMOF synthesis, (iii) characterisation of aMOFs, (iv) investigation
of local structure and defects, (v) characterising the extended structure, and (vi) establishing design
principles for aMOF synthesis. In turn, we evaluate the NC, OP, and SW approaches while exploring the
impact of composition and methodology on aMOF structure.

Nanocluster Synthesis and Formation
Before developing the NC approach, we first investigated the atomic structure and formation kinetics of the
M O nanocluster itself. Examining the kinetics of M O formation is particularly important for understanding
3 3
the OP and SW synthesis methods later. The M O nanocluster is synthesised by combining concentrated
3
solutions of a metal nitrate salt and sodium acetate, followed by heating and crystallisation [see Methods].
To study this reaction, we employed in situ pair distribution function (PDF) analysis using X‑ray total
scattering measurements in a custom flow‑cell setup (MAX IV, Sweden), where sodium acetate is injected
into the metal salt solution [Supplementary Fig. 1]. This allowed us to monitor the nanocluster formation
in real time as the reagents are combined. Specifically, the emergence of the metal–metal (M–M) distance
(~3.3 Å) in the PDF provided a direct means to track nanocluster formation in solution [Supplementary
Fig. 2].

The results revealed that the rate of nanocluster formation is highly dependent on metal chemistry. Fe O
3
forms almost instantaneously upon the addition of sodium acetate, whereas Cr O forms significantly more
3
slowly. By fitting the temporal evolution of the M–M PDF peak intensity to a simple empirical model, we
extracted pseudo‑rate constants for the reactions [Supplementary Note 1 and Supplementary Fig. 2].
The analysis showed that Fe O forms almost 18 times faster than Cr O. These differences are attributed to
3 3
the metals' electronic configurations and ligand exchange rates. Fe3+, with its high‑spin d5 configuration,
undergoes rapid ligand exchange, whereas Cr3+, with a d3 configuration, exhibits much slower ligand
exchange, the slowest among first‑row transition metals.40 These substantial differences in reaction speeds
were anticipated to have significant consequences for the formation of M‑BDC systems later on.

Next, large single crystals of the M
3
O nanoclusters were synthesSised from solutions with nominal
compositions of Fe , Fe Cr, FeCr , and Cr [see Methods].26 X‑ray fluorescence (XRF) confirmed the
3 2 2 3
bimetallic composition of the Fe Cr and FeCr nanoclusters [Supplementary Fig. 3]. Semi-quantitative
2 2
analysis suggested a slightly lower proportion of Cr than in the Enominal ratios [Supplementary Table 1].
This is unsurprising, given the significantly slower reaction raRte of Cr3+ compared to Fe3+ observed in our in
situ PDF studies. Nonetheless, single‑crystal X‑ray diffraction confirmed that the M O nanocluster structure
3
was present across all four samples [Fig. 1b and Supplementary Table 2]. This finding is consistent with
previous research showing that the structure of M O nanoclusters is minimally affected by the inclusion of
3
different metal cations.41,42

aMOF Synthesis & Optimisa tion
Extensive research has identified the synthetic conditions that optimise MOF crystallinity.43,44 Building on
this knowledge, we focused on conditions opposite to those typically used to produce highly crystalline
materials. For crystalline Fe‑BDC systems, high temperatures, long reaction times, and acidic conditions
are known to favour crystallinity, wIch is consistent with thermodynamic control.16 In contrast, to achieve
aMOFs, we hypothesised that rapid, room‑temperature syntheses under basic conditions would promote
disorder. This approach deliRberately deviates from standard practices aimed at attaining crystallinity.

To test this hypothesis, we systematically studied Fe‑BDC and Cr‑BDC, aiming to produce materials whose
diffraction patterns contained no Bragg peaks. Starting with a synthesis optimised for a highly crystalline
Fe‑BDC polymorph, we systematically varied key parameters, including temperature, reaction time, reagent
concentration, modulator (acetic acid) volume, and pH.16,26 Details of these studies are provided in
Supplementary Note 2, Supplementary Figs. 4 and 5 and Supplementary Table 3.

In summary, room‑temperature reactions using high reactant concentrations, low modulator volumes, and
the addition of base were optimal for forming aMOFs from both Fe O and Cr O nanoclusters. Navigating
3 3
the ligand exchange rates of Fe3+ and Cr3+ was crucial in determining the optimal synthesis conditions for
aMOF formation.40 The significantly slower ligand exchange kinetics of Cr3+ required reaction conditions
that facilitated sufficient metal–linker interactions to enable Cr‑BDC to form an amorphous network. At the
same time, the faster ligand exchange of Fe3+ meant that conditions had to be carefully tuned to suppress
coordination reversibility, which would otherwise favour the formation of crystalline Fe‑BDC polymorphs.
Balancing these competing factors was essential for achieving an amorphous product across different metal
compositions.

Having established suitable conditions for directly synthesising aMOFs, we synthesised our complete series
of samples across various compositions and methods. We synthesised M‑BDC using three synthetic
methods (NC, OP, and SW) and four compositions (Fe, Fe Cr, FeCr and Cr), yielding a total of 12 aMOF
2 2,
samples [see Methods]. This systematic approach allowed us to compare the influence of both composition
and synthesis methods on the resulting aMOFs. M-BDC samples were obtained as a viscous gel-like
sample upon mixing and washing. Upon drying in air, these gels formed large monolithic pieces that took
the shape of their container [Supplementary Figure 6]. Then, upon drying in an oven, macroscopic pieces
were obtained. For measurement, the monoliths were ground into a fine powder.

Preliminary characterisation of the 12 aMOF samples was performed using X‑ray total scattering and
powder diffraction measurements [Fig. 1c and Supplementary Figure 7]. The scattering patterns for all
samples exhibited broad and poorly defined regions of diffuse scattering. One sample, FeCr ‑BDC (OP),
2
contained a minor crystalline impurity phase identified as unreacted linker material [Supplementary Figure
8]. A comparison between the scattering patterns revealed that the regions of diffuse scattering were largely
consistent across all samples. However, slight variations were observed: Fe‑BDC (SW) showed slightly
more defined scattering regions, while Cr‑BDC (NC) exhibited less well‑defined regions. XRF spectroscopy
confirmed the presence of both Fe and Cr in all six mixed‑metal samples, and semi-quantitative analysis
indicated higher Fe:Cr for Fe Cr‑BDC samples than FeCr ‑BDC, as expected [Supplementary Figure 9
2 2
and Supplementary Table 4]. Scanning electron microscopy images from the three Fe-BDC samples
revealed inhomogeneous morphologies for each synthesis method [Supplementary Figure 10].

Evolved gas analysis (EGA) was conducted to evaluate the decomposition temperatures of the M‑BDC
samples [Fig. 2a & Table 1]. For comparison, EGA measurements were also performed on the Fe O and
3
Cr
3
O nanoclusters and the H
2
BDC linker [Supplementary Fig. 11]. As pSreviously reported, Cr
3
O exhibited
a higher decomposition temperature than Fe O.45,46 For H BDC, a rapid mass loss was observed at 310 °C
3 2
due to sublimation of the linker.47

The thermal decomposition of M‑BDC occurred in two distinct stages: (1) desolvation and dehydroxylation,
followed by (2) decomposition of the organic linker. The first stage accounted for 14–22% of the total mass
loss in most samples, except for Fe‑BDC (SW), which exhibited a 38% mass loss [Table 1]. We note that
stage (2) often proceeded via a complex, multi-stage decomposition route, which has previously been
observed for other highly defective aMOF systems.48

TG‑MS analysis showed that the initial mass loss in most samples was due to the loss of DMF
[Supplementary Figure 12]. This is consistent with previous reports showing that DMF can coordinate as
the terminal (axial) ligand on M O clusters.21 Cr‑BDC (NC) was an exception, in which no DMF release was
3
detected. Instead, the mass loss in this sample was attributed almost entirely to dehydroxylation
[Supplementary Figure 13]. The mass loss associated with this process (~17%) was comparable to the
loss of DMF from the other M-BDC samples, suggesting that Cr‑BDC (NC) contains a much higher degree
of hydroxylation. In this case, the terminal hydroxo/aqua ligands from the aqueous nanocluster synthesis
likely remain coordinated, as the slow ligand-exchange kinetics of Cr(III) hinder substitution by DMF upon
aMOF synthesis in DMF. This, in turn, implies a lower proportion of Cr sites bound to the organic linker,
indicating a greater degree of metal–linker defects. The second stage, corresponding to linker
decomposition, resulted in a mass loss of 41–56% in most samples [Table 1]. However, Cr‑BDC (NC)
exhibited a significantly lower mass loss of only 27%, suggesting reduced organic linker content.

The decomposition temperature of each sample was determined as the point at which the highest rate of
mass loss occurred, as identified from the derivative thermogram (DTG) [Table 1]. Across all M-BDC
samples, decomposition temperatures ranged from 288 to 447 °C, with the most significant variations
primarily attributed to differences in composition rather than the specific synthesis method used. A clear
trend was observed: as the Cr content increased, the decomposition temperature systematically decreased.
This behaviour contrasts with that of the isolated M O nanoclusters, for which Cr O exhibits greater intrinsic
3 3
thermal stability than Fe O. The decrease in decomposition temperature with increasing Cr content
3
therefore suggests a loss of structural integrity of the Cr
3
O clusteEr within the amorphous network, implying
that additional structural motifs are present beyond intact M O SBUs. Further discussion of this structural
3
instability is provided later. Additionally, some samples exhibited multiple peaks in the DTG, indicating that
decomposition occurred in distinct stages, likely due to stepwise breakdown of the framework. This
behaviour aligns with previous reports on aMOFs, furth er reinforcing the complex nature of their thermal
decomposition.48

TGA also allowed us to estimate the degree of saturation of the M‑BDC samples [Table 2]. A fully
coordinated sample (i.e., one without metaEl–linker defects) corresponds to 100% saturation. This was
achieved by assuming that the residual Lmass after decomposition corresponds solely to the metal oxide
(M O ), following the method outlined in Ref. 49 [Supplementary Note 3]. These saturation values are semi-
2 3 C
quantitative but internally self-consistent and are therefore most meaningful when considered
comparatively rather than absolutely. Additionally, the composition of each sample was also estimated
[Supplementary Table 5].

No strong correlation wAas observed between the chemical composition of M‑BDC and the degree of
saturation. However, the NC approach generally produced samples with lower degrees of saturation. This
suggests that using pre‑formed SBUs leads to a higher concentration of metal–linker defects, likely due to
ligand exchange between the capping acetate anions and the BDC linker during synthesis.

Several systems exhibited high degrees of saturation (>90%), indicating very few metal–linker defects.
Conversely, Cr‑BDC (NC) exhibited an exceptionally low degree of saturation of 31%, consistent with its
reduced linker content. These results reinforce the idea that Cr‑BDC (NC) has a significant proportion of
metal–linker defects, aligning with the greater hydroxylation observed in this sample.

We used Fourier transform infrared (FT‑IR) spectroscopy to probe the binding modes present in M-BDC
[Fig. 2b]. FT‑IR spectra were recorded in the range of 400 to 4000 cm−1 [Supplementary Fig. 14]. Key
bands at 1560 and 1380 cm−1 were observed in all samples, corresponding to the asymmetric and
symmetric stretching vibrations of the carboxylate group, respectively.41,50 These bands confirm the metal–
linker interaction, with their positions indicating that the BDC linker coordinates to the metal centres in a
bidentate fashion. However, the asymmetric stretch at 1560 cm−1 was notably broad in most samples,
suggesting a degree of linker unsaturation, consistent with the metal–linker defects inferred from TGA.51
Broadening in FT-IR spectra is also indicative of increased structural disorder, which is further reflected in
the complex and sample-dependent features observed in the fingerprint region (~1250–750 cm−1), where
variations in aromatic C–H bending and carboxylate deformation modes are expected to be particularly
sensitive to local coordination heterogeneity and disorder.

The M–O stretch, observed around 520 cm−1, was similar for Fe‑BDC and Fe₂Cr‑BDC.52 However, the peak
showed noticeable shifts, broadening, and asymmetry at higher Cr content. These effects were most
pronounced in Cr‑BDC (NC), where significant peak broadening and shifts indicate inhomogeneity in Cr–O
interactions, likely due to local structural variations. Other features in the FT‑IR spectrum of Cr‑BDC (NC)
align with the TGA findings, supporting the conclusions of increased hydroxylation and reduced linker
content in this sample. For example, the reduced intensity and presence of a prominent shoulder on the
symmetric carboxylate stretch, as well as the presence of an additional band at 2900 and 3400 cm−1, which
correspond to the C–H stretch of a methyl group and hydroxyl groups, respectively
[Supplementary Fig. 14].53 These observations may indicate that Cr‑BDC (NC) exhibits distinct structural
characteristics relative to other M‑BDC materials and may retain some acetate capping groups, further
highlighting the influence of composition and synthesis methods on aMOF structure.

Nitrogen adsorption isotherms were measured at 77 K for the three Fe-BDC members of the family
[Supplementary Figure 15]. Both Fe-BDC (OP) and (SW) were non-porous, with maximum adsorbed
values of only 8.8 and 0.3 cm3 g−1, respectively. For Fe-BDC (NC), a type III isotherm was observed with a
maximal uptake of 135.8 cm3 g−1. We note that under the drying conditions employed here, low-porosity
materials were expected to form.54

Our results indicate that we have successfully formed amorphous M‑BDC materials and have revealed
varying degrees of metal–linker defects, with trends linked to composition and synthesis method.
Collectively, these results emphasise the critical role of Cr and thEe NC approach in driving the formation of
highly defective frameworks.

Local Structure and Node Defects
A primary goal of our study was to obtain aMOFs with a well‑defined local structure. We used pair
distribution function (PDF) analysis of the total scattering data to investigate the atomic structure of the
M‑BDC materials. Visual comparison of the PDFs for all 12 samples revealed striking similarities, with key
peaks observed at ca. 1.5, 2.0, 3.0, and 3.4 Å [Fig. 3a]. Moreover, we evaluated the reproducibility of the
synthesis methods and found that almost identical PDFs were obtained in each case [Supplementary
Figure 16].

We first consider the PDFs of the nanoclusters themselves. PDFs calculated from our experimentally
derived single‑crystal structures of the M O nanocluster were nearly identical across metal chemistries,
3
highlighting the negligible impact of metal composition on the nanocluster’s core structure in the solid state
[Fig. 3b]. The positions of the characteristic peaks in the calculated PDFs closely matched those in the
experimental PDFs, corresponding to (i) M–O, (ii) O–O and (iii) M–M atomic distances.

Moving on to the experimental aMOF PDFs, the primary distinction between the M O nanocluster and the
3
M‑BDC PDFs was the relative intensity of the M–M (ca. 3.3 Å) to the M–O (2.0 Å) peaks, which varied with
sample composition. For Fe‑BDC, the relative peak intensity aligned with the expected ratio from the
nanocluster. However, as the proportion of Cr increased, the relative intensity of the M–M peak decreased,
suggesting a lower degree of condensation in the local structure, that is, fewer M–M correlations for a given
number of M–O bonds.

To explain this, we explored alternative local structural motifs, including octahedra, dimers, and chains, as
potential causes of the observed decrease in M–M intensity. The rationale for these motifs is as follows: (i)
octahedra and dimers have been noted as intermediates in the formation of the Cr O nanocluster and (ii)
3
chains are a natural choice arising from the known polymorphism in the M‑BDC crystalline systems.16,55,56
Support for the presence of such defect nodes also comes from the thermal analysis: Cr-BDC (NC) exhibits
the lowest decomposition temperature in the series, despite Cr O clusters being more thermally stable than
3
Fe O clusters in isolation. This suggests that, within the amorphous network, the Cr O motif is not fully
3 3
preserved and likely undergoes partial dissociation, consistent with incorporation of defect nodes such as
dimers or octahedra. We also note that the high degree of hydroxylation observed for Cr-BDC (NC) is
compatible with such partially dissociated motifs. Alternative defect nodes would require extensive
reorganisation or ligand redistribution, which is unlikely given the slow ligand-exchange kinetics of Cr3+.

Given that the M–M peak remains observable at all compositions, albeit at lower intensity, this rules out the
possibility of purely octahedral or dimeric structures (since neither possesses the characteristic M–M peak).
Conversely, although the chain structure contains the M–M peak, its PDF is almost identical to that of the
trimer itself, eliminating the possibility of a solely chain-based structure.

Thus, given the EGA and FT‑IR results, which indicated varying degrees of unsaturation (i.e., metal–linker
defects) in all samples, we hypothesised that the local structure consisted of trimers alongside defective
nodes such as octahedra, dimers, or chains. By calculating linear combinations of the PDF for trimers and
these defect motifs, we could simulate what different populations of node/defect-node might look like and
see if these can capture the experimentally observed variation [Fig. 3c]. The linear combination with the
chain motif failed to induce sufficient changes in the relative intensity of the M–M peak and was not
considered further. On the other hand, both the trimer–octahedra and trimer–dimer combinations were able
to augment the M–M peak successfully. The trimer–octahedra combination also showed a reduction in the
O–O peak at 4.5 Å, while the trimer–dimer combination introduced shifts in the M–O and O–O peaks and
a new peak around 3.7 Å.

This approach allowed us to estimate the proportion of defect nodes from the relative intensities of the M–
O and M–M peaks [Supplementary Note 4]. As it was challenging to definitively distinguish between
octahedral and dimeric defects (or a combination of the two), the proportion of defects was quantified for
both scenarios [Table 3]. The defect proportion varied more with composition than with the synthesis
method. The NC approach consistently yielded fewer defects among synthetic methods, while the OP
approach produced the highest defect concentrations. Mixed‑metal samples exhibited more significant
defect concentration variability than Fe‑BDC or Cr‑BDC.

To complement the PDF analysis, we also performed X -ray absorption spectroscopy at both Fe and Cr K-
edges for all M-BDC materials. In the XANES region of the spectra, all Fe and Cr features were very similar
[Supplementary Figure 17 & Supplementary Table 6]. In all Fe-containing samples, the Fe edge position
was ~7125 eV, and the expected pre-edge feature at 7113 eV was observed, both consistent with analogous
Fe-BTC materials containing Fe3+ in an octahedral environment.57,58 In all Cr-containing samples, the Cr
edge position was 6001 eV, and the expected pre-edge feature at 5990 eV was observed, both consistent
with a related, crystalline material, MIL-100 (Cr), containing Cr3+ in an octahedral environment.59 The
EXAFS region of the spectra was then transformed from energy to k-space up to 15 Å−¹, and the Fourier
transform was calculated to obtain a real-space map of the Fe and Cr scattering path lengths
[Supplementary Figures 18 & 19]. The uncorrected Fourier transform of χ(k) was dominated by a
prominent peak between 1.47 and 1.56 Å for Fe and 1.50 to 1.57 Å for Cr, corresponding to the M–O peaks.
The slight variations in this peak indicate some degree of inhomogeneity in the local structure, as commonly
expected for an aMOF, particularly given the defective nature of these systems. Similar variations are
observed in the higher-R region of the Fourier transform, again consistent with the presence of a
superposition of varying defect structures.

Ultimately, the key finding is that Cr content is the primary factor influencing the proportion of node defects
in the M‑BDC samples, with the synthesis method playing a secondary role. We previously highlighted the
slower ligand-exchange rate of Cr3+ compared to Fe3+, which directly results in slower assembly of the Cr O
3
nanocluster than Fe O. Given that aMOFs represent metastable states, these kinetic effects also strongly
3
influence the formation of Cr‑containing M‑BDC frameworks. Under kinetic control, incorporating Cr results
in a greater degree of defects as reversible ligand exchange becomes increasingly hindered, preventing
the exchange necessary to ‘heal’ misconnections. This emphasises the significant role of composition in
influencing local structure in aMOFs.

Extended Structure
While the local structure of M-BDC frameworks is primarily influenced by chemistry, both composition and
synthetic approach play significant roles in determining the extended structure. The PDFs reveal four main
clusters of overlapping peaks in the medium range (centred around 7, 9, 11, and 13 Å) and three distinct
shapes in the long range [Fig. 4a]. Most samples exhibit a broad hump around ~22 Å, while Fe-BDC (SW)
and FeCr -BDC (OP) display weak peaks, and Cr-BDC (NC) is entirely flat. Beyond 30 Å, all samples are
2
featureless, reflecting their amorphous nature.

The NC approach consistently yields the least-ordered frameworks among the three synthesis methods, as
evidenced by shorter coherence lengths (i.e., the maximum real-space distance over which atomic pair
correlations persist, reflecting the size of structurally correlated domains in a material). [Table 4]. Due to
Cr-induced defects, the coherence length decreases systematically with increasing Cr content for the NC
and SW methods. However, this trend is absent in the OP method. While NC samples exhibit the lowest
coherence lengths, neither OP nor SW consistently produces the most extended order, suggesting a more
stochastic influence of these syntheses on structural coherence. We note that coherence length and node
defect concentration are related but not directly correlated descriptors of disorder: coherence length reflects
the spatial extent of network correlations, whereas node defects describe local deviations in node geometry,
and the two can therefore vary independently.

Even in the absence of a structural model, statistical analyses of PDFs can provide robust insights. To
quantify structural similarities among the experimental PDFs, we used the Pearson correlation coefficient,
which ranges from +1 (complete correlation) to −1 (complete anti-correlation).60,61 Then, hierarchical
clustering was used to organise these similarities into a dendrogram, where closely related PDFs cluster at
shorter branch heights, that is, towards the bottom of the dendrogram, indicating greater similarity [Fig.
4b].62 This approach offers an alternative to the full Pearson correlation matrix, effectively grouping
structurally similar materials [Supplementary Fig. 20].

This analysis immediately identifies five separate groups of samples, with Cr-BDC (NC) identified as the
most dissimilar sample: (i) Cr-BDC (NC), (ii) Fe-BDC (SW), (iii) Fe-BDC (NC), Fe Cr-BDC (NC), FeCr -
2 2
BDC (NC), and Fe-BDC (OP), (iv) Cr-BDC (OP, SW) and (v) Fe Cr-BDC (OP, SW) and FeCr -BDC (OP,
2 2
SW). By plotting these five main groups of PDFs, we can see that the analysis successfully clusters similar
PDFs together [Supplementary Fig. 21]. This naturally leads to the question: How do their atomic
structures differ?

To gain insight into the underlying atomic structure, we compared the M-BDC aMOFs to known crystalline
polymorphs: MIL-101, MIL-88B, MOF-235, MIL-53, and MIL-68 [Fig. 4c].16 While we do not expect perfect
agreement between our highly disordered aMOFs and these ordered crystalline structures, these
comparisons provide a reference for interpreting the medium-range order in our samples. MIL-101, MIL-
88B, and MOF-235 are trimer-based frameworks [Fig. 5a]. MIL-101 features nanoclusters assembled into
supertetrahedral SBUs, while MIL-88B and MOF-235 have simpler networks formed simply by bridged
nanoclusters. MIL-53 and MIL-68, in contrast, consist of chains of octahedra linked by BDC, and given the
varying degrees of defect sites in M-BDC, these motifs may also be present [Fig. 5b]. Another key difference
is the relative orientation of the trimer or chains with respect to the linker: in MIL-101 and MIL-68, they adopt
a cis-configuration, whereas MIL-88B, MOF-235, and MIL-53 adopt a trans-configuration.

The calculated PDFs for these crystalline polymorphs display sharp peaks extending beyond 20 Å, in stark
contrast to the broad features of the aMOFs [Fig. 4c]. However, certain peak positions align with
experimental data. To account for the structural disorder in the aMOFs, we considered how to artificially
introduce disorder into the crystalline structures. We chose a simple approach: increasing the thermal
displacement parameter (B ), which broadens the peaks and simulates the effect of structural disorder.
iso
Subtle differences were observed in the medium-range order (5 to 15 Å), while more pronounced variations
appeared in the long range (> 15 Å). However, as most experimental features fall within the medium range,
definitive classification of the aMOFs remains challenging. Nevertheless, statistical analysis provided
further insight by calculating the Pearson correlation coefficient between each experimental PDF and the
five crystalline polymorphs.

Notably, instead of different PDF groups aligning with different polymorphs, nearly all M-BDC samples—
except Cr-BDC (NC)—correlated most strongly with MOF-235 [Supplementary Fig. 22]. This suggests a
preference for a trans-orientation of trimers relative to the linker. Cr-BDC (NC) did not correlate with any
structure, which is expected given the absence of extended structural features in its PDF. When the order
of experimental PDFs in the correlation matrix is rearranged to match the hierarchical clustering groups, a
clear trend emerges [Fig. 5c]. Each group of PDFs exhibits similar correlation values with the crystalline
polymorphs, reinforcing that the structural motifs within each group are closely related.

Rather than trying to define aMOFs as direct amorphous analogues of crystalline phases, this analysis
reveals a more nuanced picture: no single polymorph fully describes any sample, as indicated by the
highest Pearson correlation coefficient reaching only 0.76. Our findings challenge the common assumption
that aMOFs can be directly mapped onto known crystalline polymorphs. Unlike amorphous materials
derived from crystalline counterparts, where a clear structural relationship exists, these aMOFs were
synthesised directly under amorphous-favouring conditions. This raises a broader question: What defines
a unique amorphous phase? Must both short- and medium-range order align with a crystalline reference,
or must the framework topology be partially retained? This remains an open question. These results
emphasise that aMOFs cannot simply be viewed as ‘disordered versions’ of known crystalline structures.
Instead, they represent a distinct, structurally complex class of materials with unique atomic arrangements
that emerge from their specific synthesis conditions.

Two key trends emerge by examining the relationship between node defect concentration and coherence
length across the explored phases: chemistry primarily dictates defect density, while the synthesis method
strongly influences coherence length [Fig. 6]. These findings provide a fundamental understanding of how
disorder can be controlled in aMOFs and serve as the foundation for broader design principles.

aMOF Design Principles
Building on these insights, we propose three system-informed design principles for the direct synthesis of
aMOFs. Rather than being universally applicable, these principles reflect trends observed across the
materials studied here and are intended to highlight conditions that favour amorphous network formation
over crystallisation in nanocluster-derived systems. They provide a conceptual framework for rationally
targeting aMOFs and for understanding how synthesis parameters influence disorder and defect chemistry.

(1) Kinetic Bias Towards Amorphisation: In the systems studied here, fast reactions at room temperature
using high reagent concentrations under basic conditions bias formation toward amorphous networks by
suppressing ordered assembly. Under these conditions, limited ligand exchange and reduced coordination
reversibility disfavour crystallisation pathways. While similar conditions may also support crystalline MOF
formation in other systems, our results show that kinetically dominated regimes can be exploited to access
aMOFs when combined with appropriate building units, with fine-tuning of reaction rate and reversibility
remaining critical.

(2) Nanocluster Approach as a Route to Enhanced Disorder: Using pre-formed SBUs promotes higher
structural disorder compared to direct mixing of metal salts and linkers, likely due to the additional ligand
exchange and network reorganisation required during assembly. In the present system, this results in
aMOFs with shorter coherence lengths. Although nanocluster-based strategies are not yet widespread in
aMOF synthesis, our results indicate that they represent a powerful, though currently underexplored,
pathway to controllable amorphous networks.

(3) Defect Engineering via Compositional Tuning: Incorporation of more kinetically inert metals increases
defect concentration and structural disorder by limiting structural rearrangement. For example, Cr
incorporation leads to higher defect densities than in Fe-based analogues, reflecting slower ligand-
exchange kinetics. This demonstrates that metal composition can be used as a design handle to modulate
disorder and defect landscapes, with direct implications for properties such as reactivity and host–guest
interactions, though the specific outcomes will depend on the metal–ligand chemistry involved.

Taken together, these guidelines outline how kinetic control, choice of building unit, and metal chemistry
can be combined to bias synthesis toward amorphous rather than crystalline materials. They are not
intended as universal rules, but as a transferable design framework that can inform future efforts to
deliberately access and exploit disorder in aMOFs beyond conventional crystallisation-driven approaches.

In this study, we have demonstrated the direct synthesis of amorphous metal–organic frameworks (aMOFs)
using pre-formed nanoclusters and rigid organic linkers. By systematically investigating the interplay
between composition, synthesis conditions, and structural disorder, we have established design principles
that govern aMOF formation. We find that using nanoclusters as synthetic precursors enables the formation
of highly disordered frameworks. Additionally, our findings show that kinetic control, particularly fast
reactions under basic conditions at room temperature, is crucial for suppressing crystallinity and promoting
amorphous network formation.

A key outcome of our study is the role of metal composition in defect engineering. Incorporating kinetically
less labile metals, such as Cr, significantly increases defect concentration by limiting structural
rearrangement, thereby providing a powerful tool for tuning aMOF properties. Our analysis of local and
extended structure highlights that while aMOFs exhibit clear structural motifs related to their crystalline
counterparts, they cannot be directly mapped onto known crystalline polymorphs. Instead, these materials
form unique amorphous networks dictated by both synthesis conditions and composition.

Importantly, our results reinforce that aMOFs are not merely disordered versions of crystalline frameworks.
Attempts to correlate their structures with known crystalline polymorphs reveal that no single ordered
framework fully describes their atomic arrangements. Instead, aMOFs possess distinct network
architectures that demand new characterisation approaches beyond traditional crystallography. This
highlights the necessity of multi-scale structural analysis: while the choice of SBUs primarily determines
local order, the synthesis conditions and composition strongly influence the extended disorder. Our work
demonstrates the power of PDF analysis coupled with statistical clustering in resolving these complex
structural relationships.

Moving beyond the conventional crystallise–amorphise paradigm, our bottom-up study provides a
foundation for the rational design of next-generation aMOFs with controlled disorder and tailored
functionality. These principles allow for aMOF design in applications requiring structural flexibility, defect
engineering, and enhanced transport properties, marking a step toward the rational design of amorphous
framework materials.

Methods
Nanocluster synthesis
Nanocluster synthesis was performed as described in Ref. 26. M(NO ) .9H O (0.015 mol) [Table 5] was
3 3 2
dissolved in 5 mL hot water (50 °C), and CH COONa (0.03 mol) was dissolved separately in 5 mL hot water
3
(50 °C). The two solutions were mixed and stirred for 10 minutes at 50 °C. The resulting solution was
allowed to cool down to room temperature and left in air for crystallisation. After several days, single crystals
were obtained.

Table 5 M(NO3)3.9H2O synthesis ratios for nanocluster synthesis.
Nanocluster Fe(NO3)3.9H2O Cr(NO3)3.9H2O
Fe3O 0.015 mol 0.000 mol
Fe2CrO 0.010 mol 0.005 mol
FeCr2 0.005 mol 0.010 mol
Cr3O 0.000 mol 0.015 mol

Single-crystal X‑ray diffraction
Single‑crystal X‑ray diffraction data were collected using a Bruker D8 Venture diffractometer with a
PHOTON 100 CMOS detector, a MoKα high brilliance IμS X‑ray tube (λ = 0.71073 Å), a multilayer X‑ray
mirror, and the temperature kept at 100 K using an Oxford Cryo system. The structures were solved using
Olex2 with the ShelXT program using intrinsic phasing.63,64 Refinement was done using the ShelXL
refinement package with least-squares minimisation. The extinction coefficient was refined. Hydrogen
atoms were added automatically using a riding model. All other atoms were refined anisotropically.

X‑ray fluorescence spectroscopy
XRF measurements were carried out using a Malvern Panalytical Epsilon 4 XRF spectrometer equipped
with an Ag‑anode X‑ray tube operating at 50.000 kV/6 μA and a high‑resolution silicon drift detector.
Samples were prepared in 35 mm disposable plastic cups on 3.6 μm PETP polyester film. Before
measurement, the sample environment was filled with He‑gas to avoid detecting elements in the air.

M‑BDC Syntheses
All chemicals were used as received without further purification. Iron (III) nitrate nonahydrate (99%),
Chromium (III) nitrate nonahydrate (99%), terephthalic acid (98%), sodium acetate (99%), sodium
hydroxide pellets (98%), and N, N-dimethylformamide (99.8%) were all purchased from Sigma Aldrich.

One-pot (OP) M(NO ) .9H O (0.6 g, 1.5 mmol) and H BDC (0.25 g, 1.5 mmol) were dissolved in 5 mL DMF.
3 3 2 2
NaOH (0.06 g, 1.5 mmol) was slowly added, and the mixture was stirred for 5 minutes. The product was
washed three times with DMF (5 mL) and three times with 60 °C deionised water (5 mL) and dried in an
oven at 60 °C.

Stepwise (SW) M(NO ) .9H O (0.6 g, 1.5 mmol) and NaCH COO (0.37 g, 4.5 mmol) were dissolved in 5
3 3 2 3
mL DMF. H BDC (0.25 g, 1.5 mmol) was slowly added to the mixture, and the mixture was stirred until
2
dissolved. NaOH (0.06 g, 1.5 mmol) was then slowly added, and the mixture was stirred for 5 minutes. The
product was washed three times with DMF (5 mL) and three times with 60 °C deionised water (5 mL) and
dried in an oven at 60 °C.

Nanocluster (NC) M O cluster (100 mg, 0.15 mmol) and H BDC (75 mg, 0.45 mmol) were dissolved in 5
3 2
mL DMF. NaOH (36 mg, 0.9 mmol) was then slowly added, and the mixture was stirred for 5 minutes. The
product was washed three times with DMF (5 mL) and three times with 60 °C deionised water (5 mL) and
dried in an oven at 60 °C.

Powder X-ray diffraction
Finely ground M-BDC samples were deposited onto a zero-background sample holder and mounted on the
diffractometer. Measurements were performed at room temperature on a Bruker D8 Discover diffractometer
equipped with Cu Kα radiation (λ = 1.54 Å) in Bragg-Brentano geometry. Data were collected for one hour
between 12 and 70° 2θ with a step size of 0.02°, using an Fe fluorescence filter.

Scanning electron microscopy
SEM was measured with the lower electron detector of a JEOL 7800F-prime using an acceleration voltage
of 5 kV and a working distance of 10 mm. Finely ground powder of Fe-BDC was dispersed in ethanol and
drop-cast onto Si wafers before being adhered to the sample holder. The quality of the 100k-magnification
images was limited by the instrument's resolution and the buildup of electric charge on the poorly conducting
sample. This can cause particles to move during raster scanning and, in some cases, degrade the sample
during data collection.

Fourier transform infrared spectroscopy
The IR spectra were recorded using a Bruker ALPHA‑P FT‑IR spectrometer with a single-reflection ATR
module. Data were collected in transmission mode between 375 and 4000 cm−1 on a powdered sample.

Thermogravimetric and mass spectrometry measurements
Thermogravimetric (TG) analysis coupled with mass spectrometry (MS) and infrared spectroscopy (FTIR)
was conducted using a Netzsch TG 209 F1 Libra (Netzsch, Germany) apparatus, simultaneously coupled
with an Aëolos QMS403C mass spectrometer and a Bruker spectrometer. MS data acquisition was
controlled by the OPUS (Bruker, Germany) and Aëolos (Netzsch, Germany) software. The quadrupole mass
spectrometer (QMS) detected ion currents from the gaseous products over a range of 1−50 m/z. They were
collected from room temperature to 500 °C with a heating rate of 10 °C min−1 under an air atmosphere (40
mL min−1), and around 5 mg of sample was used for each measurement. To account for any balance drift
in the instrument due to increasing heat, each TG measurement was initiated with a buoyancy correction
run, conducted on an empty crucible. This correction mirrored the gas-flow and thermal-ramp conditions.
Solvent and linker mass losses were determined from the plateau region of the TGA curve. Solvent mass
loss was taken as the difference in mass between the initial plateau and the onset of linker decomposition,
identified by a sharp mass drop and significant detection of mass 44 in the MS. Linker decomposition mass
loss was determined from the mass change between the plateau before the sharp decomposition drop and
the stable plateau after decomposition.

Nitrogen adsorption measurements
Adsorption isotherms of nitrogen were measured at 77 K using a Micromeritics TriStar II PLUS instrument.
Approximately 40 mg of each sample was degassed at 60 °C for 18 h prior to measurement.

X‑ray total scattering and pair distribution function analysis
X‑ray total scattering data were collected at the DanMAX (MAX IV, Sweden) and P02.1 (DESY, Germany)
beamlines using synchrotron radiation (λ = 0.354 Å and λ = 0.207 Å, respectively). A small amount of finely
ground sample was loaded into a Kapton capillary (1 mm diameter). Capillaries were sealed and mounted
onto the instrument. Data were collected under ambient conditions for each sample, empty capillary, and
the blank instrument over the region ~0.2 < Q < ~18.4 Å−1. The raw total scattering data were corrected for
background, multiple scattering, container scattering, and Compton scattering using xPDFsuite over the
range 0.5 < Q < 17 Å−1 using an r value of 1.0 and Q of 18.4 Å−1.65
poly max,inst
In situ X-ray total scattering data were collected at the DanMAX beamline (MAX IV, Sweden) using
synchrotron radiation (λ = 0.354 Å). A custom flow setup was used to perform the in situ measurement.
Aqueous solutions of the metal salt were placed in a round-bottom flask and pumped through a fused silica
capillary (0.85 mm diameter) placed in the path of the beam using a peristaltic pump. A separate syringe
pump was then used to inject the sodium acetate solution and initiate the reaction. The reaction mixture
was heated to 50 °C using a hot plate. Data were measured in the range ~0.85 < Q < 17.5 Å−1 using a
DECTRIS PILATUS3 X CdTe 2M area detector and a time resolution of 1 s. Aqueous solutions of the metal
salt were used to facilitate background subtraction, and the PDFs were generated over the range 0.5 < Q
< 16 Å−1 using an r value of 1.0 and Q of 18 Å−1. Fitting of the M–M peak intensity was performed
poly max,inst
in Python using the SciPy package.66

X-ray absorption spectroscopy
Spectroscopic measurements were performed at beamline B18 (core EXAFS) at Diamond Light Source
(Didcot, United Kingdom).67 The synchrotron was operating with a ring energy of 3 GeV and at a current of
300 mA. The X-ray source was a bending magnet, and energy scanning was performed in continuous
QEXAFS mode using a Si(111) dual-crystal monochromator (DCM). X-ray focusing to the sample position
was achieved using Si mirrors coated with Pt, pitched at 2.4 mrad, yielding an X-ray spot size of 600 × 600
μm2 FWHM (h × v). Data were acquired in transmission mode using ion chambers and Stanford amplifiers.
Acquisition times of 131.21 s (over the Fe K-edge) and 182.46 s per spectrum (over the Cr K-edge) were
used with a 0.3 eV step size. The energy calibration of the DCM was achieved by measuring the spectra of
Fe and Cr foils. Data were processed in Athena, wherein a flat background from the pre-edge slope and a
fitted low-order polynomial to the post-edge region were both subtracted.68 The Fourier transform of the
normalised data were also performed in Athena using a k-weight of 2.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "The manuscript presents a strategy for developing amorphous metal-organic frameworks (aMOFs) using nanoclusters as building blocks combined with rigid BDC ligands. By employing three synthesis methods—nanocluster (NC), one-pot (OP), and stepwise (SW)—along with four different compositions (Fe, Fe2Cr, FeCr2, and Cr), the study systematically compares how both composition and synthesis methods influence the resulting aMOF structures using advanced characterization techniques. Overall, this work offers an interesting approach to controlling aMOF structures through synthetic conditions and precursor selection. However, addressing the following points will significantly enhance the clarity, robustness, and overall impact of the study:\n1. Prior to this work, some examples of amorphous or glassy metal-organic frameworks (MOFs) and coordination polymers have been synthesized using a \"direct synthesis\" approach, some of which also employed metal clusters as building units. It is beneficial to mention these studies in the introduction and highlight the advancements this work brings in comparison.\n2. Powder X-ray diffraction of all aMOF samples should be presented.\n3. In the introduction, the authors state, “By targeting highly disordered frameworks, we seek to maximise sample isotropy, a property that has the potential to enhance mass transport, while controlling defect concentrations can be leveraged for improved catalytic performance.” To reinforce the impact of this strategy, it is required to include experimental results related to mass transport and/or catalytic properties. This would provide stronger support for the claimed advantages of the synthesized aMOFs.\n4. The exact composition of all aMOFs synthesized in this work should be provided. Combining results from techniques such as elemental analysis and ICP would further strengthen the estimated degree of saturation obtained from thermogravimetric analysis.\n5. Please verify the assigned decomposition temperatures in Table 1. Additionally, provide an explanation or references for why the peak position in DTA is used instead of the onset temperature. For example, at 447 °C for Fe-BDC (NC), the sample is already near the final mass fraction ((cid:1)f) as defined in Supplementary Note 3 and should have long surpassed the decomposition temperature.\n6. In Figure 2B and Supplementary Figure 11, FTIR peaks should be labeled. The overlapping spectra are difficult to read. Moreover, the noticeable differences in peaks between 1250 and 750 cm⁻¹ require discussion.\n7. For all PDF results, including a structural model with labeled peaks is highly recommended, as it would help readers better follow the discussion. Additionally, highlighting the peak positions corresponding to the linkage between nanoclusters is crucial, as this would provide direct evidence of their connectivity.\n8. Given the growing interest in amorphous MOFs, I recommend including a detailed methodology for calculating and analyzing the Pearson correlation coefficient used in this work.", "reviewer_2": "The manuscript by Sapnik, Jensen et al. reports on a method for direct synthesis of amorphous metal-organic frameworks (MOFs) from rigid organic linkers and iron and chromium nanoclusters. Based on structural (x-ray total scattering) and thermal analyses, the work focuses on comparing this method to traditional methods that rely on prior crystal formation. It is an important and interesting work for opening the doors towards tailoring functionality and disorder of amorphous MOFs. I hope the authors can address the following comments before publication:\n1. In the Introduction, the authors refer to amorphous MOFs, citing refs. 5-11. However, the references refer to studies on both MOFs (and coordination polymers) in amorphous state and glass state. Amorphous solids and glasses are both non-crystalline but distinct (e.g., see the work of Gupta: https://doi.org/10.1016/0022-3093(95)00502-1). This terminology should be clarified to avoid any confusion. I believe all the studied MOFs herein are amorphous, at least no experimental evidence of glass transition temperatures has been reported.\n2. In relation to the previous comment, the “enhanced processability” of amorphous MOFs relative to crystalline MOFs is mentioned. This is certainly correct in the case for glassy MOFs, which can be formed and processed like traditional glasses due to the access to the liquid state. But is it also correct that amorphous MOFs are more processable than crystalline MOFs? Also, in the same statement, it is not clear what is meant by “mechanical stability”, please specify.\n3. The quest of addressing “what defines an amorphous counterpart to a crystalline MOF?” is interesting and relevant. While I agree that this has not yet been explored in detail for MOFs, it has been in the case of traditional glass families, such as oxides. To my knowledge, such studies go back to at least 1992 (https://doi.org/10.1524/zkri.1992.200.3-4.287), where local modifier cation arrangements (alkali-oxygen distance and coordination numbers) of silicate glasses and crystals of the same composition were compared. In such systems, there is also not always a “crystalline parent structure”. I therefore think it would be interesting to put the current study in the context of the large literature that exists in these related fields.\n4. The provided x-ray PDF data are nice and insightful but building accurate structural understanding from a single technique is challenging for amorphous materials. Therefore, I think additional structural analysis is needed to fully prove the authors claims about the local structures and better understand the chemical nature of the interactions. This could include neutron PDF, X-ray absorption spectroscopy (EXAFS, XANES), or Mössbauer spectroscopy (in case of Fe). At least for one or more selected samples, I think it is needed to corroborate the structural models with an additional structural characterization method, or alternatively atomistic simulations.\nMinor points:\n5. A reference to Supplementary Fig. 5 is missing in Supplementary Note 2. I assume the discussed results in this section are referring to the data in this figure.\n6. The blue (Cr-BDC) and purple (FeCR2-BDC) curves in Figs. 1-4 are hard to distinguish.", "reviewer_3": "The manuscript submitted by Nanna and co-workers reports the direct synthesis of amorphous metal–organic frameworks (aMOFs) from preformed nanoclusters (i.e., secondary building units, SBUs). By employing M₃O (M = Fe, Fe₂Cr, FeCr₂, Cr) nanoclusters and 1,4-benzenedicarboxylate (BDC) linkers under kinetically controlled conditions (room temperature, base-mediated), the authors synthesized a series of amorphous M-BDC materials (12 samples in total). The amorphous nature of these aMOFs (e.g., defects and structural disorder) was examined using a range of techniques, including X-ray pair distribution functon(PDF) analysis, thermogravimetric analysis (TGA) with evolved gas analysis, Fourier-transform infrared spectroscopy (FTIR), and statistical PDF analysis. Based on their findings, the authors propose three key design principles for aMOF synthesis: (1) kinetic control, (2) the use of preformed SBUs, and (3) defect engineering via compositional tuning. Overall, this manuscript is of interest and provides insight into the synthesis, characterization, and defect engineering of aMOFs. However, the manuscript would benefit from additional clarification and discussion on several critical points before it can be accepted for publication. I recommend publication after major revision, addressing the following concerns.\nMajor Comments\n1. The rapid reaction under high reagent concentrations often results in uncontrollable amorphous structures, particularly with regard to defect density and site specificity. This is one of my major concerns. I strongly recommend that the authors verify the reproducibility of the PDF and defect-related data for identical samples from different synthesis batches, or at least provide data from one or two representative samples.\n2. Pore size and surface area are critical parameters in MOF materials. Given the structural diversity and varying defect levels reported, it would be beneficial to explore whether these textural properties are influenced by the amorphous nature and defect content. Such measurements would significantly strengthen the discussion of structure-property relationships.\n3. The authors highlight that controlling defect concentrations may enhance catalytic performance. It would strengthen the manuscript to include brief experimental validation-if possible-even as a preliminary proof of concept, to demonstrate the practical utility of amorphous engineering.\n4. The manuscript would benefit from a clearer definition of what constitutes a \"defect\" in the context of these aMOFs. Are defects defined as missing linkers, incomplete coordination, local disorder, or a combination thereof? Clarifying this will improve understanding of the characterization results.\n5. Please clarify why DMF remains in the aMOFs after water washing, with Cr-BDC (NC) being the only exception. Could this retention be related to the anomalously low decomposition temperature of Cr-BDC (NC) shown in Table 1?\nMinor Comments\n1. The authors cite previous computational studies on amorphous Fe–BDC frameworks. A brief comparison of the current experimental findings with those computational results would provide useful context and strengthen the manuscript's relevance to the broader literature.\n2. There appears to be a discrepancy in how defect content is reported: TGA suggests a higher degree of metal–linker defects (i.e., lower coordination saturation) in SBU-based aMOFs, while PDF modeling indicates a lower proportion of defect nodes in the same samples.\n3. Table 2 seems to show highly random variability in the saturation degree of different samples prepared by the SW and OP methods. This also raises concerns about batch-to-batch reproducibility, which should be addressed or commented upon.\n4. The authors present three design principles for aMOF synthesis. However, these may not be universally applicable. For instance, base conditions can also be favorable for some crystalline MOF syntheses, and nanocluster-based approaches are not yet mainstream. It would be better to moderate their claims and clarify the scope of their design strategy.\n5. Many parts of the manuscript delve into highly technical discussions too quickly. Given the journal’s broad readership, the authors are encouraged to (1) define specialized terms (e.g., coherence length); (2) spell out acronyms at first mention (e.g., ORTE); and (3) cite relevant literature (e.g., when referring to different MILs). Pls double-check throughout the whole MS. Additionally, please clarify the terms “aq” and “s” used in Figure 1. What is the distinction between M₃O (aq) and M₃O (s)?\n6. The authors claim that aMOFs synthesized from SBUs exhibit greater structural disorder, while the one-pot method appears to result in higher defect levels based on Figure 6. This contradiction should be clarified.", "reviewer_4": "Dr. Sapnik et al.’s work on amorphous metal-organic frameworks (aMOFs) is a transformative contribution to a field that has long prioritized crystalline systems. By addressing the critical synthesis and control challenges of aMOFs—topics largely overlooked by the MOF community—this study fills a vital gap in materials science. The meticulous analysis of pair distribution functon(PDF) data and dynamic properties, coupled with the exploration of short-, mid-, and long-range structural order, provides unprecedented insights into these complex amorphous systems. The authors’ bold interrogation of established paradigms in MOF and hybrid materials research is both refreshing and intellectually stimulating, pushing the boundaries of how we conceptualize disorder in porous materials.\nThe question, “What defines a unique amorphous phase?” is elegantly framed and resonates as a cornerstone inquiry in modern materials science. Answering this could indeed catalyze a paradigm shift, unlocking new design principles for amorphous materials with tailored functionalities. This work’s interdisciplinary impact—spanning synthesis, characterization, and theoretical modeling—makes it a compelling candidate for Nature Communications, where it will inspire diverse researchers to rethink the role of disorder in material properties.\nMinor Comments/Curiosities:\nThe sol-gel nature of aMOFs under ambient conditions raises intriguing questions about their thermal response. Would heating induce crystallization, or might it instead yield robust amorphous monoliths? Clarifying this could deepen understanding of their stability and potential applications.\nGiven the porosity-centric focus of MOF research, briefly addressing the porosity profile of these aMOFs (e.g., surface area, pore accessibility) would further highlight their functional relevance, even in amorphous states.\nWhile the PDF analysis in Figure 1 is insightful, the overlapping traces slightly obscure detail. Could the authors adjust line weights or colors for clarity? Additionally, the pronounced spike at ~2 Å for FeCr₂-BDC is intriguing—is this attributable to distinct metal-linker bonding or local structural motifs?\nThese minor points only underscore the richness of this elegant study. Warmest congratulations to the authors on this exceptional contribution!"}, "author_response": "RESPONSE TO REVIEWERS’ COMMENTS\nReviewer #1 (Remarks to the Author):\nThe manuscript presents a strategy for developing amorphous metal-organic frameworks (aMOFs) using\nnanoclusters as building blocks combined with rigid BDC ligands. By employing three synthesis methods—\nnanocluster (NC), one-pot (OP), and stepwise (SW)—along with four different compositions (Fe, Fe2Cr, FeCr2,\nand Cr), the study systematically compares how both composition and synthesis methods influence the resulting\naMOF structures using advanced characterization techniques. Overall, this work offers an interesting approach to\ncontrolling aMOF structures through synthetic conditions and precursor selection. However, addressing the\nfollowing points will significantly enhance the clarity, robustness, and overall impact of the study:\n1. Prior to this work, some examples of amorphous or glassy metal-organic frameworks (MOFs) and coordination\npolymers have been synthesized using a \"direct synthesis\" approach, some of which also employed metal clusters\nas building units. It is beneficial to mention these studies in the introduction and highlight the advancements this\nwork brings in comparison.\nWe have now expanded on the discussion of direct syntheses in the introduction and highlighted recent\nnanocluster-based work.\nWhile a small number of studies have demonstrated that direct synthesis of aMOFs is possible by kinetically\nsuppressing crystallisation, these examples remain comparatively rare and system-specific, underscoring that\ngeneralisable design rules for the rational, reproducible synthesis of aMOFs are still largely lacking.12,15\nWhile direct synthesis of aMOFs from nanoclusters has been demonstrated using highly flexible linkers to\nkinetically suppress crystallisation, the combination of nanoclusters with rigid organic linkers to generate aMOFs\nremains essentially unexplored.30\nAdditional references:\n12. Lin, Z., Richardson, J. J., Zhou, J. & Caruso, F. Direct synthesis of amorphous coordination polymers and\nmetal–organic frameworks. Nat Rev Chem7, 273–286 (2023).\n15. Molavi, H. Amorphous metal-organic frameworks: From synthesis strategies to emerging applications. Journal\nof Environmental Chemical Engineering13, 117890 (2025).\n30. Ma, N. et al. Systematic design and functionalisation of amorphous zirconium metal–organic frameworks.\nChem. Sci.15, 17562–17570 (2024).\n2. Powder X-ray diffraction of all aMOF samples should be presented.\nWe have now included powder X-ray diffraction data for all 12 M-BDC samples, confirming their amorphous nature.\nPreliminary characterisation of the 12 aMOF samples was performed using Xray total scattering and powder\ndiffraction measurements [Fig.1c and Supplementary Fig. 6].\n-\nAdded to methods:\nPowder X-ray diffraction\nFinely ground M-BDC samples were deposited onto a zero-background sample holder and mounted on the\ndiffractometer. Measurements were performed at room temperature on a Bruker D8 Discover diffractometer\nequipped with Cu Kα radiation (λ = 1.54 Å) in Bragg-Brentano geometry. Data were collected for one hour between\n12 and 70° 2θ with a step size of 0.02°, using an Fe fluorescence filter.\nSupplementary Figure 7 Powder X-ray diffraction patterns from all 12 of the M-BDC materials.\n3. In the introduction, the authors state, “By targeting highly disordered frameworks, we seek to maximise sample\nisotropy, a property that has the potential to enhance mass transport, while controlling defect concentrations can\nbe leveraged for improved catalytic performance.” To reinforce the impact of this strategy, it is required to include\nexperimental results related to mass transport and/or catalytic properties. This would provide stronger support for\nthe claimed advantages of the synthesized aMOFs.\nUnfortunately, performing mass transport/catalytic studies is beyond the scope of this work, which focuses on\nsynthesis and structure. Nevertheless, these enhanced properties are well-described in the literature, and we have\ntherefore added additional references here to support our statement.\nAdditional references:\n15. Molavi, H. Amorphous metal-organic frameworks: From synthesis strategies to emerging applications.\nJournal of Environmental Chemical Engineering 13, 117890 (2025).\n31. Yu, Z., Tang, L., Ma, N., Horike, S. & Chen, W. Recent progress of amorphous and glassy coordination\npolymers. Coordination Chemistry Reviews 469, 214646 (2022).\n4. The exact composition of all aMOFs synthesized in this work should be provided. Combining results from\ntechniques such as elemental analysis and ICP would further strengthen the estimated degree of saturation\nobtained from thermogravimetric analysis.\nThe compositions of the aMOFs have now been included.\nSupplementary Table 1 Estimated compositions for each sample. The composition of an aMOF of the M-BDC\ntype is M3O(OH2)(OH)(C8O4H4)x and in the case where there are no linker/cluster defects, x = 3.\nFe- Fe- Fe- Fe2Cr- Fe2Cr- Fe2Cr- FeCr2- FeCr2- FeCr2- Cr- Cr- Cr-\nBDC BDC BDC BDC BDC BDC BDC BDC BDC BDC BDC BDC\n(NC) (SW) (OP) (NC) (SW) (OP) (NC) (SW) (OP) (NC) (SW) (OP)\nx 2.3 3.2 2.4 2.2 4 3.1 1.6 2.6 2.7 0.8 2.8 2.3\n5. Please verify the assigned decomposition temperatures in Table 1. Additionally, provide an explanation or\nreferences for why the peak position in DTA is used instead of the onset temperature. For example, at 447 °C for\nFe-BDC (NC), the sample is already near the final mass fraction ( f) as defined in Supplementary Note 3 and\nshould have long surpassed the decomposition temperature.\n𝑚\nWe agree that the decomposition temperature could have also been described using the onset temperature.\nHowever, we chose to report the derivative of the TGA as the decomposition of these samples is often complex,\nproceeding via a multi-stage decomposition route. We have previously reported this phenomena for a related\nsystem in Ref. 48 and found that the “stepped” nature of the decomposition increased with defect concentration.\nBecause of this, it is difficult to directly compare the onset temperatures when one or more decomposition steps\nare present. Consequently, we chose to report the derivative of the TGA as a more consistent value for the different\nmaterials.\nWe note that stage (2) often proceeded via a complex, multi-stage decomposition route, which has previously been\nobserved for other highly defective aMOF systems.[48]\n48. Sapnik, A. F. et al. Stepwise collapse of a giant pore metal–organic framework. Dalton Trans.50, 5011–5022\n(2021).\n6. In Figure 2B and Supplementary Figure 11, FTIR peaks should be labelled. The overlapping spectra are difficult\nto read. Moreover, the noticeable differences in peaks between 1250 and 750 cm ¹ require discussion.\nWe have now added labels to highlight the key bands in both figures and also⁻ reduced the line thickness, as\nsuggested by the other reviewers.\nFigure 2b:\nSupplementary Figure 14:\nWe have also expanded our discussion to explicitly discuss the fingerprint region as follows:\nBroadening in FT-IR spectra is also indicative of increased structural disorder, which is further reflected in the\ncomplex and sample-dependent features observed in the fingerprint region (~1250–750 cm−¹), where variations in\naromatic C–H bending and carboxylate deformation modes are expected to be particularly sensitive to local\ncoordination heterogeneity and disorder.\n7. For all PDF results, including a structural model with labelled peaks is highly recommended, as it would help\nreaders better follow the discussion. Additionally, highlighting the peak positions corresponding to the linkage\nbetween nanoclusters is crucial, as this would provide direct evidence of their connectivity.\nThe key M–M and M–O correlations have now been highlighted in Figure 1c and the peak corresponding to inter-\ncluster correlations has been highlighted in Figure 4a.\nFigure 1c:\nFigure 4a:\n8. Given the growing interest in amorphous MOFs, I recommend including a detailed methodology for calculating\nand analyzing the Pearson correlation coefficient used in this work.\nWe have now added details regarding the calculation of the Pearson correlation coefficient to the Methods section\nof the manuscript.\nFitting of the M–M peak intensity and calculation of the Pearson correlation coefficient were both performed in\nPython using the SciPy package.65\nAdditional reference:\n65. Virtanen, P. et al. SciPy 1.0: fundamental algorithms for scientific computing in Python. Nat Methods17, 261–\n272 (2020).\nReviewer #2 (Remarks to the Author):\nThe manuscript by Sapnik, Jensen et al. reports on a method for direct synthesis of amorphous metal-organic\nframeworks (MOFs) from rigid organic linkers and iron and chromium nanoclusters. Based on structural (x-ray total\nscattering) and thermal analyses, the work focuses on comparing this method to traditional methods that rely on\nprior crystal formation. It is an important and interesting work for opening the doors towards tailoring functionality\nand disorder of amorphous MOFs. I hope the authors can address the following comments before publication.\n1. In the Introduction, the authors refer to amorphous MOFs, citing refs. 5-11. However, the references refer to\nstudies on both MOFs (and coordination polymers) in amorphous state and glass state. Amorphous solids and\nglasses are both non-crystalline but distinct (e.g., see the work of Gupta:https://doi.org/10.1016/0022-\n3093(95)00502-1). This terminology should be clarified to avoid any confusion. I believe all the studied MOFs\nherein are amorphous, at least no experimental evidence of glass transition temperatures has been reported.\nWe agree this should be clarified; the references given do indeed refer broadly to non-crystalline MOFs, which\nincludes glassy MOFs. We have now commented that our focus here is specifically amorphous MOFs, a type of\nnon-crystalline MOF which does not possess a glass transition.\nDepending on the system and conditions, such treatments can also yield glassy MOFs that exhibit a glass\ntransition, representing one subset within the broader family of non-crystalline MOF materials.10\nIn this work, we challenge the ‘crystallise–amorphise’ paradigm by developing a nanocluster-based approach for\nthe direct synthesis of aMOFs, specifically non-crystalline MOFs without a glass transition.\n2. In relation to the previous comment, the “enhanced processability” of amorphous MOFs relative to crystalline\nMOFs is mentioned. This is certainly correct in the case for glassy MOFs, which can be formed and processed like\ntraditional glasses due to the access to the liquid state. But is it also correct that amorphous MOFs are more\nprocessable than crystalline MOFs? Also, in the same statement, it is not clear what is meant by “mechanical\nstability”, please specify.\nYes, the aMOFs studied here are accessed through a gel state that can be moulded into arbitrary shapes, similar\nto the casting of glasses. Upon drying in air, these gels form monolithic bodies that retain the shape of their\ncontainer, and subsequent oven drying yields dense xerogel monoliths. Because of this intrinsic gel-to-monolith\ntransition, the aMOFs are naturally obtained as single, centimetre-scale monolithic pieces rather than powders,\nunlike the vast majority of crystalline MOFs.\nTo clarify what we mean by enhanced processability and mechanical stability: crystalline MOFs typically require\npelletisation or incorporation into a binder to form monoliths. Pelletised MOFs often suffer from low mechanical\nstability and poor cohesion, while binder-based monoliths require favourable MOF–binder interfacial interactions\nand inevitably dilute the active MOF content. In contrast, our aMOF gels self-assemble into binder-free monoliths\nwith sufficient mechanical integrity to be handled, cut, and processed directly. We now clarify this distinction in the\ntext and provide images of the resulting monoliths in the Supplementary Information.\nIn contrast, amorphous MOFs (aMOFs) are gaining attention for their enhanced processability, catalytic properties,\nand greater mechanical integrity when formed as monoliths.5–11\nM-BDC samples were obtained as a viscous gel-like sample upon mixing and washing. Upon drying in air, these\ngels formed large monolithic pieces, taking the shape of their container [Supplementary Figure 6]. Then, upon\ndrying in an oven, macroscopic, like pieces were obtained. For measurement purposes, the monoliths were then\nground into a fine powder.\nSupplementary Figure 1 Example photographs of Fe-BDC once it has been (a) dried and (b) removed from the\nsynthesis container, demonstrating its monolithic form and retention of the container shape. (c) Shows the same\nsample after it has been lightly ground.\n3. The quest of addressing “what defines an amorphous counterpart to a crystalline MOF?” is interesting and\nrelevant. While I agree that this has not yet been explored in detail for MOFs, it has been in the case of traditional\nglass families, such as oxides. To my knowledge, such studies go back to at least 1992\n(https://doi.org/10.1524/zkri.1992.200.3-4.287), where local modifier cation arrangements (alkali-oxygen distance\nand coordination numbers) of silicate glasses and crystals of the same composition were compared. In such\nsystems, there is also not always a “crystalline parent structure”. I therefore think it would be interesting to put the\ncurrent study in the context of the large literature that exists in these related fields.\nWe have amended the introduction here to reflect the wide-ranging nature of this challenge, incorporating relevant\nreferences to oxide glasses, metallic glasses and amorphous semi-conductors.\nThis raises the broader question: what defines the amorphous counterpart to a crystalline material? It is a question\nthat has long been posed across families of important materials, such as oxide glasses, metallic glasses, and\namorphous semiconductors.23–25 This issue becomes especially significant when amorphous materials are\nsynthesised directly, and therefore, no obvious parent crystalline material is present. Addressing this question\nrequires a systematic approach to direct aMOF synthesis.\nAdditional references:\n23. Zachariasen, W. H. The atomic arrangement in glass. J. Am. Chem. Soc.54, 3841–3851 (1932).\n24. Chen, M. A brief overview of bulk metallic glasses. NPG Asia Mater3, 82–90 (2011).\n25. Saraswat, S., Tomar, V. K., Deolia, V. K. & Mehta, N. A brief review of amorphous semiconductors. Materials\nToday: Proceedings (2023) DOI:10.1016/j.matpr.2023.03.455.\n4. The provided x-ray PDF data are nice and insightful but building accurate structural understanding from a single\ntechnique is challenging for amorphous materials. Therefore, I think additional structural analysis is needed to fully\nprove the authors claims about the local structures and better understand the chemical nature of the interactions.\nThis could include neutron PDF, X-ray absorption spectroscopy (EXAFS, XANES), or Mössbauer spectroscopy (in\ncase of Fe). At least for one or more selected samples, I think it is needed to corroborate the structural models with\nan additional structural characterization method, or alternatively atomistic simulations.\nWe have now performed synchrotron X-ray absorption spectroscopy (XANES and EXAFS) for all 12 materials at\nboth the Fe and Cr K-edges. The data corroborate the expected oxidation (+3) for all samples and the octahedral\ncoordination about both metal centres. Moreover, we see close agreement between samples at the different edges,\nsuggesting consistency between the local structures in the different materials.\nWe have added the following text and figures to the manuscript:\nTo complement the PDF analysis, we also performed X-ray absorption spectroscopy at both Fe and Cr K-edges\nfor all M-BDC materials. In the XANES region of the spectra, all Fe and Cr features were very similar\n[Supplementary Figure 17 & Supplementary Table 6]. In all Fe-containing samples, the Fe edge position was\n~7125 eV and the expected pre-edge feature at 7113 eV was observed, both consistent with analogous Fe-BTC\nmaterials containing Fe3+ in an octahedral environment.57,58 In all Cr-containing samples, the Cr edge position was\n6001 eV and the expected pre-edge feature at 5990 eV was observed, both consistent with a related, crystalline\nmaterial, MIL-100 (Cr), containing Cr3+ in an octahedral environment.59 The EXAFS region of the spectra was then\ntransformed from energy to k-space up to 15 Å−¹ and the Fourier transform was calculated to obtain a real-space\nmap of the Fe and Cr scattering path lengths [Supplementary Figures 18 & 19]. The uncorrected Fourier transform\nof χ(k) was dominated by a prominent peak between 1.47 and 1.56 Å for Fe and 1.50 to 1.57 Å for Cr, corresponding\nto the M–O peaks. The slight variations in this peak indicate some degree of inhomogeneity in the local structure,\nas commonly expected for an aMOF, particularly given the defective nature of these systems. Similar variations\nare observed in the higher-R region of the Fourier transform, again consistent with a superposition of varying defect\nstructures being present.\nAdditional references:\n57. Sapnik, A. F. et al. Transient intermediate in the formation of an amorphous metal–organic framework. Soft\nMatter20, 2338–2347 (2024).\n58. Sciortino, L., Alessi, A., Messina, F., Buscarino, G. & Gelardi, F. M. Structure of the FeBTC Metal–Organic\nFramework: A Model Based on the Local Environment Study. J. Phys. Chem. C119, 7826–7830 (2015).\n59. Zhang, J.-Y. et al. Adsorption of Uranyl ions on Amine-functionalization of MIL-101(Cr) Nanoparticles by a\nFacile Coordination-based Post-synthetic strategy and X-ray Absorption Spectroscopy Studies. Sci Rep5, 13514\n(2015).\nSupplementary Figure 7 XANES spectra obtained from the Fe (left) and Cr (right) K-edges for all 12 members of\nthe M-BDC family.\nSupplementary Table 6 Edge position, determined from the derivative of the normalised energy spectra, for Fe\nand Cr, for all 12 members of the M-BDC family.\nFe- Fe- Fe- Fe2Cr- Fe2Cr- Fe2Cr- FeCr2- FeCr2- FeCr2- Cr- Cr- Cr-\nBDC BDC BDC BDC BDC BDC BDC BDC BDC BDC BDC BDC\n(NC) (SW) (OP) (NC) (SW) (OP) (NC) (SW) (OP) (NC) (SW) (OP)\nFe-\nedge 7125 7125 7125 7125 7125 7124 7125 7125 7125 – – –\n(eV)\nCr-\nedge – – – 6001 6001 6001 6001 6001 6001 6001 6001 6001\n(eV)\nSupplementary Figure 18 (a)k-space and (b) corresponding uncorrected Fourier transform of the EXAFS data\nobtained from the Fe K-edge. Features observed below the Fe–O path length (~1.5 Å) are artefacts from the\nbackground.\nSupplementary Figure 19 (a)k-space and (b) corresponding uncorrected Fourier transform of the EXAFS data\nobtained from the Cr K-edge. Features observed below the Cr–O path length (~1.5 Å) are artefacts from the\nbackground.\nAdded to methods:\nX-ray absorption spectroscopy\nSpectroscopic measurements were performed at beamline B18 (core EXAFS) at Diamond Light Source (Didcot,\nUnited Kingdom).67 The synchrotron was operating with a ring energy of 3 GeV and at a current of 300 mA. The X-\nray source was a bending magnet and energy scanning was done in continuous QEXAFS mode with a Si(111) dual\ncrystal monochromator (DCM). X-ray focussing to the sample position was achieved with Si mirrors coated with Pt\npitched at 2.4 mrad, to an X-ray spot size of 600 X 600 μm2 FWHM (h x v). Data were acquired in transmission\nmode using ion chambers and Stanford amplifiers. Acquisition times of 131.21 s (over the Fe K-edge) and 182.46\ns per spectrum (over the Cr K-edge) were used with a 0.3 eV step size. The energy calibration of the DCM was\nachieved by measuring the spectrum of a Fe/Cr foil respectively. Data were processed in Athena wherein a flat\nbackground from the pre-edge slope and a fitted low-order polynomial to the post-edge region were both\nsubtracted.68 The Fourier transform of the normalised data were also performed in Athena using a k-weight of 2.\nAdditional references:\n67. Dent, A. J. et al. B18: A core XAS spectroscopy beamline for Diamond. J. Phys.: Conf. Ser. 190, 012039\n(2009).\n68. Ravel, B. & Newville, M. ATHENA and ARTEMIS Interactive Graphical Data Analysis using IFEFFIT.\nPhysica Scripta 1007 (2005) doi:10.1238/Physica.Topical.115a01007.\nIn addition, we have also included scanning electron microscopy images for the three Fe-BDC samples to highlight\ntheir irregular morphology.\nScanning electron microscopy images from the three Fe-BDC samples revealed inhomogeneous morphologies for\neach synthesis method [Supplementary Figure 10].\nAdded to methods:\nScanning electron microscopy\nSEM was measured with the lower electron detector of a JEOL 7800F-prime using an acceleration voltage of 5 kV\nand a working distance of 10 mm. Finely ground powder of Fe-BDC was dispersed in ethanol and drop-cast onto\nSi wafers before being adhered to the sample holder. The quality of the 100k magnification images was limited by\nthe resolution of the instrument as well as the build-up of electric charge on the poorly conducting sample. This can\ncause the particles to move during raster scanning, and in some cases the sample degrades during data collection.\nSupplementary Figure 10 Scanning electron microscopy images for Fe-BDC (a) OP, (b) SW, and (c) NC\nsynthesis methods.\nMinor points:\n5. A reference to Supplementary Fig. 5 is missing in Supplementary Note 2. I assume the discussed results in this\nsection are referring to the data in this figure.\nWe have added the reference to Supplementary Figure 5 in Supplementary Note 2.\n6. The blue (Cr-BDC) and purple (FeCR2-BDC) curves in Figs. 1-4 are hard to distinguish.\nAlong with the comment from Reviewer 4, we have now adjusted the line style throughout the manuscript.\nReviewer #3 (Remarks to the Author):\nThe manuscript submitted by Nanna and co-workers reports the direct synthesis of amorphous metal–organic\nframeworks (aMOFs) from preformed nanoclusters (i.e., secondary building units, SBUs). By employing M O (M =\nFe, Fe Cr, FeCr , Cr) nanoclusters and 1,4-benzenedicarboxylate (BDC) linkers under kinetically controlled\nconditions (room temperature, base-mediated), the authors synthesized a series of amorphous M-BDC m₃aterials\n(12 sam₂ples in tot₂al). The amorphous nature of these aMOFs (e.g., defects and structural disorder) was examined\nusing a range of techniques, including X-ray pair distribution function (PDF) analysis, thermogravimetric analysis\n(TGA) with evolved gas analysis, Fourier-transform infrared spectroscopy (FTIR), and statistical PDF analysis.\nBased on their findings, the authors propose three key design principles for aMOF synthesis: (1) kinetic control, (2)\nthe use of preformed SBUs, and (3) defect engineering via compositional tuning. Overall, this manuscript is of\ninterest and provides insight into the synthesis, characterization, and defect engineering of aMOFs. However, the\nmanuscript would benefit from additional clarification and discussion on several critical points before it can be\naccepted for publication. I recommend publication after major revision, addressing the following concerns.\nMajor Comments\n1. The rapid reaction under high reagent concentrations often results in uncontrollable amorphous structures,\nparticularly with regard to defect density and site specificity. This is one of my major concerns. I strongly recommend\nthat the authors verify the reproducibility of the PDF and defect-related data for identical samples from different\nsynthesis batches or at least provide data from one or two representative samples.\nWe agree that direct synthesis approaches may seem to occur in an often uncontrollable manner. To verify the\nreproducibility of the syntheses, we synthesised additional batches of the Fe-BDC OP, SW and NC samples,\nperformed synchrotron total scattering measurements and PDF analysis. We found excellent agreement, with\nalmost identical PDFs obtained for each of the three sample types.\nWe have added the following to the manuscript:\nMoreover, we evaluated the reproducibility of the synthesis methods and found almost identical PDFs were\nobtained in each case [Supplementary Figure 16].\nSupplementary Figure 2 Evaluation of the reproducibility for the three synthesis methods (OP, SW and NC) of\nFe-BDC. Each PDF has been offset for clarity\nAdded to acknowledgments:\nWe acknowledge the European Synchrotron Radiation Facility (ESRF) for provision of synchrotron radiation\nfacilities and Momentum Transfer for facilitating the measurements. Jakub Drnec is thanked for assistance and\nsupport in using beamline ID31. The measurement setup was developed with funding from the European Union’s\nHorizon 2020 research and innovation program under the STREAMLINE project (grant agreement ID 870313).\n2. Pore size and surface area are critical parameters in MOF materials. Given the structural diversity and varying\ndefect levels reported, it would be beneficial to explore whether these textural properties are influenced by the\namorphous nature and defect content. Such measurements would significantly strengthen the discussion of\nstructure-property relationships.\nIn our (and others) experiences, the textural properties of aMOFs obtained via the gel state are incredibly sensitive\nto the drying conditions used. For example, in Fe-BTC (a related aMOF system), the surface area could range from\nessentially non-porous to highly micro- and macro-porous with a BET surface area of around 1620 m2 g−1 by varying\nthe gel washing and drying procedures, see Kaskel et al., Chem. Commun., 6056–6058 (2009).\nTypically, samples dried in a standard oven tend to be less porous due to the capillary forces causing a collapse of\nthe gel network. While samples dried using supercritical CO2 drying equipment tend to avoid such collapse and\nretain their porosity. In this work (and several of our previous works), we have used the former approach due to\nlack of access to sc-CO2 drying equipment but nonetheless this allows us to study the atomic structure of the\nmaterials. We emphasise that low porosity does not preclude aMOFs from being functional in other applications.\nGiven we used our standard oven drying procedures in this work, we anticipated that the aMOFs would possess\nlow porosity. To confirm this and address the reviewers point, we have now measured N2 sorption isotherms for the\nthree Fe-BDC samples. We have now included these data and their discussion in the Results section.\nNitrogen adsorption isotherms were measured at 77 K for the three Fe-BDC members of the family\n[Supplementary Figure 15]. Both Fe-BDC (OP) and (SW) were non-porous, with maximum adsorbed values of\nonly 8.8 and 0.3 cm3 g−1, respectively. For Fe-BDC (NC), a type III isotherm was observed with a maximal uptake\nof 135.8 cm3 g−1. We note that, under the drying conditions employed here it was expected that low porosity\nmaterials would form.54\nSupplementary Figure 5 Adsorption (closed) and desorption (open) isotherms of N2 at 77 K for the three Fe-\nBDC members.\nAdditional reference:\n54. Lohe, M. R., Rose, M. & Kaskel, S. Metal–organic framework (MOF) aerogels with high micro- and\nmacroporosity. Chem. Commun. 6056–6058 (2009).\nAdded to Methods:\nNitrogen adsorption measurements\nAdsorption isotherms of nitrogen were measured at 77 K using a Micromeritics TriStar II PLUS instrument.\nApproximately 40 mg of each sample was degassed at 60 °C for 18 h prior to measurement.\n3. The authors highlight that controlling defect concentrations may enhance catalytic performance. It would\nstrengthen the manuscript to include brief experimental validation-if possible-even as a preliminary proof of\nconcept, to demonstrate the practical utility of amorphous engineering.\nWe agree that experimental validation of catalytic performance would be a valuable addition. However, such studies\nare currently beyond the scope of this work, which is specifically focused on providing detailed insights into\nsynthesis–structure relationships and defect formation in amorphous MOFs. That said, numerous studies on\nanalogous aMOF systems—as well as the broader field of defect engineering—have demonstrated that controlled\ndefect concentrations and amorphisation can lead to enhanced catalytic performance. To better contextualise our\ndiscussion and support this point, we have now added additional references to representative studies in these\nareas.\nBy targeting highly disordered frameworks, we seek to maximise sample isotropy, a property that has the potential\nto enhance mass transport, while controlling defect concentrations can be leveraged for improved catalytic\nperformance.15,31\nAdditional references:\n15. Molavi, H. Amorphous metal-organic frameworks: From synthesis strategies to emerging applications.\nJournal of Environmental Chemical Engineering 13, 117890 (2025).\n31. Yu, Z., Tang, L., Ma, N., Horike, S. & Chen, W. Recent progress of amorphous and glassy coordination\npolymers. Coordination Chemistry Reviews 469, 214646 (2022).\n4. The manuscript would benefit from a clearer definition of what constitutes a \"defect\" in the context of these\naMOFs. Are defects defined as missing linkers, incomplete coordination, local disorder, or a combination thereof?\nClarifying this will improve understanding of the characterization results.\nThe use of the word “defect” can indeed be confusing. In the excellent review by Fang et al. [Angew. Chem. Int.\nEd., 54, 7234–7254 (2015)], a huge variety of types of defects are presented: missing linker defects, missing node\ndefects, node distortions, dangling linkers, dislocations, and heterogeneity, amongst many others.\nIn this manuscript, we identify two main types of defect: metal–linker defects (where the metal centres and linkers\nare not fully coordinated to each other) and node defects (where metal centres are present in any non-trimer\ngeometry). These are identified by TGA and PDF analysis, respectively. We have used the term “saturation degree”\nsynonymously with metal–linker defect.\nIn response to this comment, we have carefully reviewed the manuscript to ensure that the work “defect” is not\nused on its own, without qualification (i.e., ensuring we explicitly refer to the defects as metal–linker defects or node\ndefects).\nWe have also included our definition for saturation degree at first use:\nTGA also allowed us to estimate the degree of saturation of the MBDC samples [Table 2]. A fully coordinated\nsample (i.e., one without metal–linker defects) corresponds to 100% saturation.\n-\n5. Please clarify why DMF remains in the aMOFs after water washing, with Cr-BDC (NC) being the only exception.\nCould this retention be related to the anomalously low decomposition temperature of Cr-BDC (NC) shown in Table\n1?\nA plausible explanation for the retention of DMF after water washing is coordination of DMF to the metal clusters.\nIn the M3O nanocluster, each metal centre typically carries an additional terminal ligand that is neither the bridging\ncarboxylate nor the central μ3-oxo group. This site is often depicted as aqua or hydroxo, but coordination by DMF\nat this position has been reported (e.g., Ref. 21). Such DMF ligation during synthesis would not necessarily be fully\ndisplaced during washing, particularly if ligand exchange kinetics are slow, leading to DMF remaining bound in the\nfinal aMOF (Reviewer Figure 1).\nThis mechanism also provides a consistent explanation for why Cr-BDC (NC) is the only sample in which no DMF\nis detected. The nanoclusters are formed in water, so the terminal ligands are expected to be aqua/hydroxo, and\nthe slow ligand exchange rate of Cr(III) would hinder subsequent coordination of DMF during aMOF assembly in\nDMF. Consequently, Cr-BDC (NC) uniquely lacks coordinated DMF.\nFinally, it is plausible that coordinated DMF contributes to thermal stabilisation of the Fe-containing aMOFs. The\nabsence of such coordination in Cr-BDC (NC) could therefore help explain its anomalously low decomposition\ntemperature relative to the other materials.\nWe have now addressed this in the manuscript:\nTGMS analysis showed that the initial mass loss in most samples was due to the loss of DMF [Supplementary\nFig.12]. This is consistent with previous reports showing that DMF can coordinate as the terminal (axial) ligand on\nM3O- clusters.21\nIn this case, the terminal hydroxo/aqua ligands from the aqueous nanocluster synthesis likely remain coordinated,\nas the slow ligand-exchange kinetics of Cr(III) hinder substitution by DMF upon aMOF synthesis in DMF.\nReview Figure 1 Schematic of terminal ligation by DMF to the M3O nanocluster\nMinor Comments\n1. The authors cite previous computational studies on amorphous Fe–BDC frameworks. A brief comparison of the\ncurrent experimental findings with those computational results would provide useful context and strengthen the\nmanuscript's relevance to the broader literature.\nThe previous computational studies were on the related Fe-BTC (not BDC) systems. We have now clarified that\nour present work is on the Fe-BDC family.\nTo address this, we focus on the related FeBDC family, a system known for its extensive crystalline\npolymorphism.16\n-\n2. There appears to be a discrepancy in how defect content is reported: TGA suggests a higher degree of metal–\nlinker defects (i.e., lower coordination saturation) in SBU-based aMOFs, while PDF modelling indicates a lower\nproportion of defect nodes in the same samples.\nWe thank the reviewer for highlighting this point. In response, we have clarified the terminology used throughout\nthe manuscript, as the term “defect” is employed in the literature to describe two distinct concepts: (i) metal–linker\ndefects, referring to reduced coordination saturation (missing linkers), and (ii) node defects, referring to structural\nmotifs where entire metal nodes deviate from the ideal connectivity. In our work, PDF modelling quantifies node\ndefects, whereas TGA provides information related to linker saturation. These quantities are therefore not expected\nto correlate directly. It is entirely possible for a sample to exhibit a low fraction of defect nodes (i.e., most nodes\nremain intact), while still possessing substantial metal–linker vacancies that reduce its saturation degree. We have\nnow clarified this distinction in the text to avoid confusion.\nWe also thank the reviewer for drawing attention to our discussion of the dimer model. Upon revision, we realised\nthat our argument should be based solely on the trend in decomposition temperature, rather than the saturation\ndegree inferred from TGA. The key observation is that the decomposition temperature decreases with increasing\nCr content—opposite to what would be expected if the material were composed exclusively of M3O clusters,\ngiven the higher intrinsic thermal stability of Cr3O clusters relative to Fe3O clusters. This anomalous trend therefore\nsupports the presence of an additional structural motif—consistent with the mixed cluster/dimer model derived from\nour PDF analysis—rather than contradicting the defect analysis.\nTGA section:\nThis behaviour contrasts with that of the isolated M3O nanoclusters, for which Cr3O exhibits greater intrinsic thermal\nstability than Fe3O. The decrease in decomposition temperature with increasing Cr content therefore suggests a\nloss of structural integrity of the Cr3O cluster within the amorphous network, implying that additional structural\nmotifs are present beyond intact M3O SBUs.\nPDF section:\nSupport for the presence of such defect nodes also comes from the thermal analysis: Cr-BDC (NC) exhibits the\nlowest decomposition temperature in the series, despite Cr3O clusters being more thermally stable than Fe3O\nclusters in isolation. This suggests that, within the amorphous network, the Cr3O motif is not fully preserved and\nlikely undergoes partial dissociation, consistent with incorporation of defect nodes such as dimers or octahedra.\nWe also note that the high degree of hydroxylation observed for Cr-BDC (NC) is compatible with such partially\ndissociated motifs. Alternative defect nodes would require extensive reorganisation or ligand redistribution, which\nis unlikely given the slow ligand-exchange kinetics of Cr3+.\n3. Table 2 seems to show highly random variability in the saturation degree of different samples prepared by the\nSW and OP methods. This also raises concerns about batch-to-batch reproducibility, which should be addressed\nor commented upon.\nThe primary purpose of the saturation analysis in Table 2 is to highlight systematic differences between synthesis\nstrategies, rather than to imply precise or absolute quantification of metal–linker defects for individual samples. The\nkey and robust trend is that the NC synthesis consistently leads to a lower saturation degrees, indicative of a higher\ndensity of metal–linker defects, relative to samples prepared via the SW and OP routes.\nThe apparent variability among SW- and OP-derived samples reflects a complex interplay between synthesis\nconditions (e.g., precursor chemistry, reaction kinetics, and solvent history) and metal identity, which is expected\nfor highly defective, amorphous systems. Importantly, the saturation values are derived from TGA using the method\ndescribed in Ref. 49, which relies on assumptions about final decomposition products and does not account for all\npossible structural motifs or heterogeneities. As such, these values should be interpreted semi-quantitatively, and\nwe have now clarified this limitation explicitly in the manuscript.\nRegarding batch-to-back reproducibility, we refer to point 1 where additional syntheses were performed, and almost\nidentical PDFs were obtained. Therefore, we are confident in the reproducibility of the syntheses.\nWe have addressed this as follows:\nTGA also allowed us to estimate the degree of saturation of the MBDC samples [Table 2]. A fully coordinated\nsample (i.e., one without metal–linker defects) corresponds to 100% saturation. This was achieved by assuming\nthat the residual mass after decomposition corresponds solely to th-e metal oxide (M2O3), following the method\noutlined in Ref. 49 [Supplementary Note 3]. These saturation values are semi-quantitative but internally self-\nconsistent and are therefore most meaningful when considered comparatively rather than absolutely.\n4. The authors present three design principles for aMOF synthesis. However, these may not be universally\napplicable. For instance, base conditions can also be favorable for some crystalline MOF syntheses, and\nnanocluster-based approaches are not yet mainstream. It would be better to moderate their claims and clarify the\nscope of their design strategy.\nWe have taken this consideration into account and re-written our design principles section to highlight that they\nmay not necessarily be universal across different aMOF chemistries:\naMOF Design Principles\nBuilding on these insights, we propose three system-informed design principles for the direct synthesis of aMOFs.\nRather than being universally applicable, these principles reflect trends observed across the materials studied here\nand are intended to highlight conditions that favour amorphous network formation over crystallisation in\nnanocluster-derived systems. They provide a conceptual framework for rationally targeting aMOFs and for\nunderstanding how synthesis parameters influence disorder and defect chemistry.\n(1) Kinetic Bias Towards Amorphisation: In the systems studied here, fast reactions at room temperature using\nhigh reagent concentrations under basic conditions bias formation toward amorphous networks by suppressing\nordered assembly. Under these conditions, limited ligand exchange and reduced coordination reversibility disfavour\ncrystallisation pathways. While similar conditions may also support crystalline MOF formation in other systems, our\nresults show that kinetically dominated regimes can be exploited to access aMOFs when combined with\nappropriate building units, with fine-tuning of reaction rate and reversibility remaining critical.\n(2) Nanocluster Approach as a Route to Enhanced Disorder: Using pre-formed SBUs promotes higher\nstructural disorder compared to direct mixing of metal salts and linkers, likely due to the additional ligand exchange\nand network reorganisation required during assembly. In the present system, this results in aMOFs with shorter\ncoherence lengths. Although nanocluster-based strategies are not yet widespread in aMOF synthesis, our results\nindicate that they represent a powerful, though currently underexplored, pathway to controllable amorphous\nnetworks.\n(3) Defect Engineering via Compositional Tuning: Incorporation of more kinetically inert metals increases defect\nconcentration and structural disorder by limiting structural rearrangement. For example, Cr incorporation leads to\nhigher defect densities relative to Fe-based analogues, reflecting slower ligand exchange kinetics. This\ndemonstrates that metal composition can be used as a design handle to modulate disorder and defect landscapes,\nwith direct implications for properties such as reactivity and host–guest interactions, though the specific outcomes\nwill depend on the metal–ligand chemistry involved.\nTaken together, these guidelines outline how kinetic control, choice of building unit, and metal chemistry can be\ncombined to bias synthesis toward amorphous rather than crystalline materials. They are not intended as universal\nrules, but as a transferable design framework that can inform future efforts to deliberately access and exploit\ndisorder in aMOFs beyond conventional crystallisation-driven approaches.\n5. Many parts of the manuscript delve into highly technical discussions too quickly. Given the journal’s broad\nreadership, the authors are encouraged to (1) define specialized terms (e.g., coherence length); (2) spell out\nacronyms at first mention (e.g., ORTE); and (3) cite relevant literature (e.g., when referring to different MILs). Pls\ndouble-check throughout the whole MS. Additionally, please clarify the terms “aq” and “s” used in Figure 1. What\nis the distinction between M O (aq) and M O (s)?\nWe agree with the reviewer₃ here and hav₃e carefully revised the manuscript to clearly explain all technical terms\nused throughout.\nThe NC approach consistently yields the least-ordered frameworks among the three synthesis methods, as\nevidenced by shorter coherence lengths (i.e., the maximum real-space distance over which atomic pair correlations\npersist, reflecting the size of structurally correlated domains in a material).\nThe term ‘ORTEP’ replaced with ‘thermal ellipsoid plot’.\nAdditional references have been added for each of the MIL-materials:\nBy reacting an Fe salt with 1,4benzenedicarboxylic acid (H BDC), multiple crystalline polymorphs can be formed,\nincluding MIL 101 (MIL; Material Institute Lavoisier), MIL88, MIL 68, MIL 53, MIL 235, and Fe(BDC)(DMF).16–21\n- ₂\n16. Bara-, D. et al. Exploring and expanding the Fe--terephth-alate m-etal–orga-nic framework phase space by\ncoordination and oxidation modulation. Materials Horizons8, 3377–3386 (2021).\n17. Férey, G. et al. A Chromium Terephthalate-Based Solid with Unusually Large Pore Volumes and Surface\nArea. Science309, 2040–2042 (2005).\n18. Serre, C., Millange, F., Surblé, S. & Férey, G. A Route to the Synthesis of Trivalent Transition-Metal Porous\nCarboxylates with Trimeric Secondary Building Units. Angew. Chem.43, 6285–6289 (2004).\n19. Whitfield, T. R., Wang, X., Liu, L. & Jacobson, A. J. Metal-organic frameworks based on iron oxide\noctahedral chains connected by benzenedicarboxylate dianions. Solid State Sciences7, 1096–1103 (2005).\n20. Fateeva, A. et al. Synthesis, Structure, Characterization, and Redox Properties of the Porous MIL-68(Fe)\nSolid. European Journal of Inorganic Chemistry2010, 3789–3794 (2010).\n21. Sudik, A. C., Côté, A. P. & Yaghi, O. M. Metal-Organic Frameworks Based on Trigonal Prismatic Building\nBlocks and the New “acs” Topology. Inorg. Chem.44, 2998–3000 (2005).\nWe have updated the caption to Figure 1 explaining the difference between (aq) and (s):\nThe M3O nanocluster is first synthesised in aqueous solution (aq) and then left to form a crystalline solid (s).\nWe have also added an additional reference for the hierarchical clustering analysis.\nThen, hierarchical clustering was used to organise these similarities into a dendrogram, where closely related PDFs\ncluster at shorter branch heights, that is, towards the bottom of the dendrogram, indicating greater similarity [Fig.\n4b].62\nAdditional reference:\n62. Nielsen, F. Hierarchical Clustering. in Introduction to HPC with MPI for Data Science 195–211 (Springer\nInternational Publishing, Cham, 2016). doi:10.1007/978-3-319-21903-5_8.\n6. The authors claim that aMOFs synthesized from SBUs exhibit greater structural disorder, while the one-pot\nmethod appears to result in higher defect levels based on Figure 6. This contradiction should be clarified.\nThis apparent contradiction arises from imprecise language, which we have now corrected. In this work, the NC\napproach results in a shorter coherence length, which we synonymously refer to as increased “structural disorder.”\nIn contrast, the SW and OP methods typically yield a higher proportion of node defects. These observations are\nnot contradictory: coherence length reflects the spatial extent of correlated order across the network, whereas node\ndefects describe deviations in local node geometry. While the two are related, they are not directly correlated, and\nit is therefore possible to observe high or low coherence lengths in materials with either high or low concentrations\nof node defects.\nWe note that coherence length and node defect concentration are related but not directly correlated descriptors of\ndisorder: coherence length reflects the spatial extent of network correlations, whereas node defects describe local\ndeviations in node geometry, and the two can therefore vary independently.\nReviewer #4 (Remarks to the Author):\nDr. Sapnik et al.’s work on amorphous metal-organic frameworks (aMOFs) is a transformative contribution to a field\nthat has long prioritized crystalline systems. By addressing the critical synthesis and control challenges of aMOFs—\ntopics largely overlooked by the MOF community—this study fills a vital gap in materials science. The meticulous\nanalysis of pair distribution functon(PDF) data and dynamic properties, coupled with the exploration of short-, mid-\nand long-range structural order, provides unprecedented insights into these complex amorphous systems. The\nauthors’ bold interrogation of established paradigms in MOF and hybrid materials research is both refreshing and\nintellectually stimulating, pushing the boundaries of how we conceptualize disorder in porous materials.\nThe question, “What defines a unique amorphous phase?” is elegantly framed and resonates as a cornerstone\ninquiry in modern materials science. Answering this could indeed catalyze a paradigm shift, unlocking new design\nprinciples for amorphous materials with tailored functionalities. This work’s interdisciplinary impact—spanning\nsynthesis, characterization, and theoretical modeling—makes it a compelling candidate for Nature\nCommunications, where it will inspire diverse researchers to rethink the role of disorder in material properties.\nMinor Comments/Curiosities:\nThe sol-gel nature of aMOFs under ambient conditions raises intriguing questions about their thermal response.\nWould heating induce crystallization, or might it instead yield robust amorphous monoliths? Clarifying this could\ndeepen understanding of their stability and potential applications.\nSimilar to Review 2’s second point, we have now discussed that M-BDC is obtained as a gel-like sample, which\ncan be air dried to form large monolithic pieces which take the shape of their container. Upon oven-drying, “glassy”\nlike monoliths were obtained. These are then ground to a fine powder for subsequent characterisation.\nAs for heating induced crystallisation, we can only be sure that heating to 60°C (the temperature used in this study),\ndoes not lead to crystallisation. Higher temperatures (i.e. more similar to traditional crystalline MOF solvothermal\nsyntheses) may cause crystallisation to occur. This has previously been observed for the related Fe-BTC, where\nheating the powder in a high-humidity atmosphere induced crystallisation (Ind. Eng. Chem. Res. 2019, 58, 19,\n7801–7807).\nGiven the porosity-centric focus of MOF research, briefly addressing the porosity profile of these aMOFs (e.g.,\nsurface area, pore accessibility) would further highlight their functional relevance, even in amorphous states.\nWe have now included these measurements and a discussion on the porosity of these samples – see Reviewer\n3’s second point.\nWhile the PDF analysis in Figure 1 is insightful, the overlapping traces slightly obscure detail. Could the authors\nadjust line weights or colors for clarity? Additionally, the pronounced spike at ~2 Å for FeCr -BDC is intriguing—is\nthis attributable to distinct metal-linker bonding or local structural motifs?\n₂\nWe have now decreased the line weight throughout the manuscript for clarity.\nThe peak at ~2Å−1 was identified as a small amount of unreacted linker, we have now addressed this in the text.\nOne sample, FeCr2 BDC (OP), contained a minor crystalline impurity phase identified as unreacted linker material\n[SupplementaryFig.8].\n-\nThese minor points only underscore the richness of this elegant study. Warmest congratulations to the authors on\nthis exceptional contribution!"}, {"round": 2, "reviewer_comments": {"reviewer_1": "The authors have conducted additional experiments and provided further analysis, which has improved the clarity of the discussion. Overall, they have addressed most of the comments raised in the previous round. I have no further concerns regarding this interesting work, and I extend my congratulations to the authors.", "reviewer_2": "The authors have adequately addressed all of my previous comments. This has become a very nice paper, congratulations.", "reviewer_3": "This manuscript has been significantly improved in response to the reviewers’ comments. All of my previous concerns have been satisfactorily addressed, and the additional analyses and clarifications substantially strengthen the work. I therefore recommend publication of the manuscript in its current form.", "reviewer_4": "The authors have addressed the comments in the previous revision. Now the paper can be accepted for publication."}, "author_response": ""}]

Decision letter