Ultrahigh strength magnesium via solidification of nanocolloid

Domain: neuroscience
Published: 2026-04-10
Editorial outcome: major_revision

Abstract
We report a simple, scalable route to produce ultrahigh-strength magnesium (Mg) via solidification of a colloidal solution containing nanoscale niobium carbide (NbC) particles suspended in liquid magnesium (Mg(l)). A single-atom-level investigation reveals that NbC exhibits spontaneous wetting with molten Mg, driven by the formation of an ordered layer of Mg atoms strongly bonded to the carbon atoms on the NbC {001} surface. This creates Mg-coated NbC (Mg@NbC) particles in liquid Mg and is referred to as Mg(l)-Mg@NbC nanocolloid. This unique and spontaneous wetting behaviour enables uniform nanoparticle dispersion in the molten Mg without external fields, and in the solidified Mg matrix without the need for thermomechanical processing. The resulting NbC dispersoids act as coherent, hard reinforcement phases, significantly strengthening the Mg matrix. As a result, the Mg-NbC material exhibits ultrahigh tensile strength and stiffness, surpassing those of all previously reported Mg alloy systems.

Introduction
Mg has a low density of 1.78 g·cm-3 and a high damping capacity with a loss factor ~11.8×10-3 (~10 times higher than Al and Steel), and thus offers tremendous potential in achieving lightweight vehicles with improved fuel economy and reduced CO emission.
2
However, the yield strengths of the most used automotive Mg alloys are in the range of 100-200 MPa, which is just nearly half that of Al alloys and 1/4 that of high strength steels; thus, the industry is not fully able to utilise their lightweight potential.
Among the major alloy strengthening mechanisms, the uniform dispersion of nanoscale secondary phase precipitates is particularly effective1. Alloying additions of rare earth (RE) elements in Mg alloys results in a significant strength improvement due to the effective precipitation of RE-containing phases upon heat treatment2, 3. Despite the promise of this approach, the high cost and limited availability of RE elements have historically hindered the widespread use of Mg-RE alloys in industry, with their potential only recently being actively explored4, 5.
As another strategy, fine-grained magnesium produced through thermomechanical processing is typically combined with precipitation hardening to achieve high strength6. However, the resulting microstructures often exhibit limited thermal stability, which compromises their performance under demanding service conditions.
Ceramic particles with high shear moduli, as an alternative approach to the in-situ formed precipitates, can serve as the non-deforming particles that favour the Orowan strengthening mechanism7. Their strengthening effect depends mainly on particle size and number density. This approach does not require specially tailored alloy compositions or complex thermomechanical processing routes. Reported data on various metallic systems highlight that in addition to ceramic particles contributing to strengthening8-12, the resulting nano-sized grain structure of the metallic matrix significantly improves the overall mechanical strength. For example, in Mg alloys, a multi-step external field treatment on molten metal followed by severe plastic deformation were conducted to disperse SiC nanoparticles13. In molybdenum alloys, the in-situ reaction of the precursor compounds14 was utilised to introduce the nano-sized ceramic particles and then a rolling process was conducted to disperse the nanoparticles and refine the Mo grain size. These ceramic phases are thermally stable and can effectively retain their strengthening effect at elevated temperatures. This is the case of Al incorporated with graphene-coated MgO nanoparticles9, via the powder metallurgy route, which showed improved strength and creep resistance at temperatures as high as 500°C. Similarly, non-equilibrium solidification in additive manufacturing was applied to incorporate La O nanoparticles in a medium entropy
2 3
alloy matrix and demonstrated further enhancement of creep property at 1093°C15
compared with the pure alloy parent phase.
We report here an ultrahigh tensile strength of ~680 MPa in Mg produced by solidification of a Mg colloid containing fine scale NbC dispersions. The favourable wetting characteristics of the NbC phase with molten magnesium enabled the processing of such high strength material. We investigate the fundamental mechanism for the atomic interaction between Mg and NbC using aberration-corrected scanning transmission electron microscopy (STEM) and electron energy loss spectroscopy (EELS). We further perform density functional theory (DFT) calculations to investigate the interfacial interactions and atomic arrangements at the liquid Mg/NbC (Mg(l)/NbC) interfaces, with an emphasis on the non-polar NbC {001} surface that terminates most produced NbC nanoparticles. This study reveals that the formation of an ordered layer of Mg atoms, strongly bonded to the carbon atoms on the NbC {001} surface, results in Mg-coated NbC (Mg@NbC) colloid, which plays a critical role in the observed favourable wetting characteristics. During solidification of the Mg(l)-Mg@NbC colloid, the Mg solidification front spontaneously engulfs the NbC dispersoids into the growing Mg grains. In the resulting solidified alloys, these embedded dispersoids can effectively contribute to strengthening, similar to the role of coherent precipitates in monolithic Mg alloys.

Results
Mg-NbC material by solidification of colloidal solution
Historically, the preparation of stable high-temperature colloidal solutions via the combination of molten metal and ceramic particulate dispersions has not been possible, due to the poor wettability of the ceramic particles by the molten metal16. We have identified the metal-ceramic (Mg-NbC) system as a promising candidate to realise this task, owing to its favourable wetting characteristics that enable the production of a high-temperature nanocolloidal solution consisting of NbC nano-dispersions in molten Mg using a simple, practical and scalable method.
For producing nanocolloidal solutions, we first compressed NbC powders with particle size of 15 nm (NbC ), 287 nm (NbC ) and 906 nm (NbC ) (see Methods)
nano submicron micron
into bulk pellets and then infiltrated molten Mg into these compressed pellets at ambient pressure (Fig. 1). The colloidal solution was then solidified using a conventional casting process to obtain the Mg-NbC material that consists of uniform NbC nanodispersoids in the Mg matrix (Supplementary Figs. 1, 2 for the case of NbC particles). During
nano
infiltration into the compressed powder pellet, molten Mg wetted the nanoparticles causing them to detach from each other and form a stable nanocolloid. This preparation method significantly simplifies the incorporation of ex-situ ceramic nanoparticles into the molten metal with a uniform dispersion. Similarly, uniform dispersions can also be seen in Mg-NbC and Mg-NbC materials prepared using this approach
submicron micron
(Supplementary Figs. 3-6). It is worth noting that, to our knowledge, the favourable wetting behaviour observed in the Mg-NbC system has not been reported in structural engineering metal/ex-situ ceramic systems with nano-sized ceramic particles. To exploit this favourable wetting behaviour and to better emulate industrial processing routes for producing larger volumes of Mg-NbC materials, the infiltrated Mg-NbC pellet was
submicron
introduced into a molten Mg bath (Supplementary Fig. 7). This process successfully yielded bulk Mg containing 1 vol.% NbC particles.
submicron
Ultrahigh tensile strength in Mg-NbC materials
The measured tensile strength of Mg-NbC materials (Fig. 2a) prepared by solidification of Mg(l)-Mg@NbC colloid is over one order of magnitude higher than pure Mg. For Mg-
NbC and Mg-NbC materials, the measured strengths are 527 MPa (~10 times
nano submicron
higher) and 678 MPa (~14 times higher), respectively. Even with such high load of ceramic particles, the tested Mg-NbCsamples exhibited a ductile fracture morphology (Supplementary Fig. 8). To further evaluate the influence of particles with various size and volume fraction, the experimental data were compared with theoretical predictions in Fig. 2b. Orowan strengthening7, 13 and load-transfer strengthening17 are considered as the dominant strengthening mechanisms (see Methods). The solid-solution strengthening is absent in these materials as the matrix is pure Mg. In the absence of grain boundaries in Mg-NbC and Mg-NbC materials (Supplementary Table 1),
nano submicron
grain boundary strengthening is also excluded. However, due to polycrystalline nature of macro-tensile samples (both Mg-1vol.%NbC and Mg-13vol.%NbC ), the
submicorn micron
grain boundary strengthening is considered (see Methods). For comparison, reported yield strength improvement (∆𝜎 ) for a range of materials containing ex-situ phase
𝑌𝑆
additions (Supplementary Table 2) are compiled and shown in the inset of Fig. 2b alongside the Mg-NbC and Mg-NbC samples. As shown, the Mg-NbC system
nano submicron
exhibits significantly higher ∆𝜎
𝑌𝑆
values compared to Mg, Al, Ti, and Fe based metal matrix composites (MMCs).
The indentation test results (Fig. 2c) show a high hardness for Mg-NbC (2.59
submicron
GPa), exceeding that of 316L stainless steel (2.45 GPa), thought to be due to the presence of hard NbC dispersions in the metal matrix. Higher specific elastic modulus and hardness for all Mg-NbC samples over other engineering alloys are also achieved (Supplementary Fig. 9).
The mechanical properties of a range of metallic materials and MMCs, measured using the micro-tensile method and reported in the literature, are tabulated in Supplementary Table 3. The product of specific ultimate tensile strength and elongation (which is an indicator for specific toughness) is plotted against the specific elastic modulus in Fig. 2d. Mg-NbC exhibits a unique combination of specific elastic modulus (~32
nano
MN·m·kg-1) and specific toughness indicator (~2500 kN·m·%·kg-1). This demonstrates the advantages of this ex-situ particle strengthening strategy for simultaneous enhancement of specific stiffness and toughness at room temperature. While these metrics clearly highlight the performance advantages of Mg-NbC, the broader applicability of this approach for advanced structural applications will ultimately depend on its high-temperature mechanical response and microstructural stability;
comprehensive investigations in this regard are currently ongoing and will be reported in a future dedicated publication.
Mg/NbC interfaces at atomic scale
To understand the wetting behaviour of NbC by molten Mg, as well as the origin of the uniform nanoparticle dispersion and the resulting mechanical performance observed in the Mg-NbC system, we now turn to a detailed investigation of the Mg/NbC interface. First, we present atomic-scale experimental evidence using scanning transmission electron microscopy (STEM) combined with electron energy-loss spectroscopy (EELS), which reveals the presence of a thin, ordered layer of Mg atoms on the surface of the NbC particles. Next, we support and extend these observations through first-principles density functional theory (DFT) simulations, which provide insights into the atomic bonding and interfacial energetics responsible for this unique surface structure.
As the introduction of nanoparticles into molten metal is the most challenging, and their strengthening effectiveness is superior among the three Mg-NbC materials synthesised and studied here, the Mg-NbC material has been selected for in-depth interfacial
nano
analysis. We employed dedicated STEM imaging combined with EELS core-loss mapping to simultaneously resolve the Mg/NbC interfacial structure and its chemical
{001}
composition. The images in Fig. 3a, b provide complementary contrast sensitive to the interfacial atomic structure29, 30, while the corresponding compositional distribution is further resolved by EELS elemental mapping in Fig. 3c. We can readily recognise interfacial atomic layers of different characteristics to the NbC substrate, which are confirmed to be rich in Mg. Predominant atomic features in each layer can be distinguished by the relative compositions across the interface (Fig. 3c, d). Notably, this interfacial multilayer complexion is epitaxially coherent with the NbC {001} plane and each Mg atomic column in the terminating Mg layer is seen to coordinate with a C atomic column immediately beneath it. This interfacial Mg atomic arrangement is referred to as C -like arrangement. As the distance away from the NbC {001} surface
NbC
increases, the Mg atoms arrangement deviates from the C -like configuration and
NbC
adopts a hexagonal close packed (HCP) structure (Supplementary Fig. 10).
First-principles DFT simulation results for the liquid Mg (Mg(l)) and NbC solid substrate system are shown in Fig.4. A snapshot of the atomic arrangements of the Mg(l)/NbC
{001}
interface (viewed along the <100> direction) equilibrated at 1000 K is presented in
NbC
Fig. 4a. The corresponding atomic arrangement along the <110> direction is
NbC
provided for completeness in Supplementary Fig. 11. Strikingly, the simulated interfacial structure shows feature similar to those observed in the post-mortem Mg-NbC
nano
material after solidification (Fig. 3). The Nb and C atoms in the substrate are well-ordered and the substrate remains solid at 1000K. The Mg atoms in the melt adjacent to the substrate (terminating Mg atoms) exhibit a layering phenomenon with a high degree of solid-like atomic ordering31 as shown in the atomic density profile (Supplementary Fig. 12). This ordering is quantified using the in-plane ordering coefficient, which is calculated as the average of local order parameters of the atoms within the individual layer at the interface32. Our calculations show that the Mg coating layer has an average in-plane ordering coefficient of 0.45, approximately half the value observed for the adjacent solid NbC layer. Each surface C atom is coordinated with one terminating Mg atom. Those terminating Mg atoms exhibit ordered structures similar to that of C atoms in face centred cubic NbC (C -like arrangement). The C-Mg interatomic distances
NbC
vary with typical lengths between 2.0 Å to 2.4 Å (Supplementary Fig. 13). The in-plane ordering coefficient rapidly decreases with increasing distance from the NbC substrate to the Mg atomic layers considered. The Mg atoms within the second layer have weak long-range ordering; this region is more liquid-like and correspondingly the in-plane ordering coefficient is just 0.04 (Fig. 4b).
The calculated net charge distribution at the Mg(l)/NbC interface is shown in Fig. 4c.
{001}
The C atoms at the NbC surface layer exhibit the same valence as in the bulk (-1.56 e-
/C). Meanwhile, surface Nb atoms lose 1.37 e-/Nb on an average, which is slightly less than that in the bulk (1.56 e-/Nb). The Mg atoms lose 0.2 e-/Mg on average, which
terminal
indicates a moderate charge transfer to the coordinating C atoms beneath. The iso-surfaces of electron density distribution at the Mg(l)/NbC interfacial area in Fig. 4d
{001}
show a high electron density (blue shade) within the NbC substrate. The Mg atoms that are bonded to C have low electron density around them, suggesting an ionic nature33 of the bonds which promoted the localised C
NbC
-like arrangement of Mg atoms adjacent to the NbC substrate. In contrast, the Mg atoms away from the interface consists of a free-electron cloud and ‘naked’ Mg ions.
Overall, the DFT calculations confirm the formation of an ordered layer of Mg atoms being strongly bonded to the surface-terminating C atoms, forming a Mg-coated NbC surface (Mg@NbC).
The combined experimental observations using atomic-resolution STEM-EELS in the as-solidified Mg-NbC sample (Fig.3) with supporting DFT calculations (Fig. 4)
nano
unravel a C -like arrangement of the ordered Mg layer on the NbC substrate. This
NbC
suggests that Mg atoms spontaneously adhere to the surface of the NbC phase, as shown in Supplementary Figure 12, driven by chemical attraction between NbC and Mg(l). Such interfacial adherence results in enhanced macroscopic wetting. During solidification, the C -like arrangement of the terminal Mg layer remains stable due to
NbC
chemical bonding between Mg and C atoms on NbC particles. A typical hcp structure of Mg forms beyond this ordered Mg multilayer, as observed in the solidified sample (Supplementary Fig. 10).
Stable Mg-NbC nanocolloid at high temperature
To achieve a stable nanocolloid in a molten magnesium medium (1000 K) with NbC nanoparticles of 15 nm diameter, the following conditions must be satisfied:
(i) The thermal energy associated with nanoparticle Brownian motion (E ) must
b
exceed the van der Waals attraction energy (W ) between particles,
vdWaals
ensuring dispersion rather than aggregation.
(ii) The energy barrier that prevents nanoparticle contact and sintering (W )
barrier
must be greater than the Brownian motion energy (E ), so that random
b
collisions do not lead to irreversible bonding or clustering.
Thermal energy for nanoparticle dispersion: E is calculated to be 13.8 zJ (E =kT, where
b b
k is the Boltzmann constant and T is the processing temperature of 1000 K).
Van der Waals attraction: The van der Waals interaction between two naked or coated
NbC nanoparticles in molten Mg can be estimated by the equation34, 35:
𝑊 = − 𝐴 𝑁𝑏𝐶−𝑀𝑔(𝑙)−𝑁𝑏𝐶∙ 𝑅𝑁𝑏𝐶 (1)
𝑣𝑑𝑊𝑎𝑎𝑙𝑠
6𝐷 2
2
𝐴 ≈ (√𝐴 −√𝐴 ) (2)
𝑁𝑏𝐶−𝑀𝑔(𝑙)−𝑁𝑏𝐶 𝑁𝑏𝐶 𝑀𝑔(𝑙)
where D, the minimum distance between two nanoparticles, is 0.2 nm (atomic layer of liquid Mg36). R , the average radius of NbC nanoparticles, is 7.5 nm. A and A
NbC Mg(l) NbC
are the Hamaker constants of molten Mg and naked NbC for the van der Waals interaction. A is 206zJ for molten Mg37 and A is calculated to be 222 zJ using
Mg(l) NbC
Lifshitz theory38, 39. Therefore, the absolute value of the maximum van der Waals attraction (W ) between two naked NbC nanoparticles in liquid Mg is estimated to
vdWaals
be 0.94 zJ, which is less than the thermal energy for Brownian motion.
In the case of Mg-coated NbC particles (Mg@NbC), Prieve and Russel38 have demonstrated that coating can mask the substrate particle and alter the Hamaker constant of the substrate particles towards the value of the coating phase (here Mg phase) when the distance is very close (D→0). Thus, the value of A in
NbC-Mg(l)-NbC
Equation (2) will be decreased significantly between Mg@NbC particles in the molten Mg medium. Therefore, for Mg coated N bC particles in the molten Mg media, the Brownian motion dominates, and the particles will be unlikely to aggregate.
Energy barrier preventing naked nanoparticle contacting: At high temperatures,
nanoparticles may sinter together upon contact, driven by a reduction in interfacial energy. In a Mg(l)-Mg@NbC colloid, as shown in Fig. 3a, a stable C -like ordered Mg
NbC
layer forms facilitated by ionic bonding between Mg and C on the outermost NbC {001} surface. For NbC particles to make contact and sinter within the molten Mg, the Mg-C bonds at the Mg(l)/NbC interface must be broken. The bonding energy of Mg-C is
nano
estimated using MgC in a rock salt-type structure40 due to its similar Mg-C bond lengths of 2.0-2.5 Å. The bonding energy41 is defined as ΔE = E - n·E , where E
(Mg-C) (Mg-Mg) (Mg-C)
and E represent the bond energies of a single Mg-C bond in MgC and a single
(Mg-Mg)
Mg-Mg bond in hcp Mg, respectively; n is the quantity of lost Mg-Mg bonds due to the formation of Mg-C bond for the Mg atoms adjacent to the NbC substrate. The calculated average interfacial bonding energy for Mg-C is -1.54 J·m-².
The debonding work (W ) required for NbC particles to come into contact and
debonding nano
sinter is given by W = S∙𝜎 , where S is the effective area of interaction
debonding 𝑀𝑔−𝐶
(calculated as S = RD with R = 7.5 nm, and D = 0.2 nm36) and 𝜎 is the interfacial
0 0 𝑀𝑔−𝐶
bonding energy, -1.54 J·m-². The resulting debonding work is calculated to be 7.25×103
zJ, which is approximately three orders of magnitude higher than the thermal energy
(E =13.8 zJ). Consequently, the nanoparticles are unlikely to contact and sinter with one
b
another under these conditions.
In brief, once naked NbC particles are introduced into liquid Mg, the formation of a Mg-coated layer on NbC leads to a stable Mg(l)-Mg@NbC nanocolloidal solution. Colloid stability for larger particles with influence of gravity is detailed in Supplementary Note 1.
NbC particle engulfment by the Mg solidification front
In typical solidification processes, ceramic particles (especially at the submicron or nanoscale) are often rejected at the solid/liquid interface due to interfacial energy considerations. This leads to particle segregation and clustering along grain boundaries, which can severely degrade the mechanical performance of the solidified metal. In contrast, as shown in Supplementary Figs. 2b, 4c, 6c, the Mg-NbC system surprisingly exhibited uniform particle dispersions with a wide range of particle sizes (from nanoscale to microscale) in the entire solidified Mg matrix, indicating a favourable interfacial interaction. This characteristic is critical for maintaining the uniform dispersion of nanoparticles initially achieved in the nanocolloidal solution and is a key factor in achieving the observed ultrahigh mechanical performance. Kaptay42 has proposed an interfacial criterion for spontaneous and forced engulfment of ceramic particles by an advancing solid/liquid interface. The sign of the interfacial force acting between a ceramic particle (c) and a solidification front (s) through the thin layer of a liquid (l) is determined by the sign of the quantity ∆𝜎 (Δσ = σ - σ  (0.08 + 1.22  cosΘ )
𝑐𝑙𝑠 cls cv lv clv
where σ is the surface energy of the ceramic, σ is the surface tension of molten metal
cv lv
and Θ is the contact angle of the molten metal on the ceramic substrate). The
clv
interfacial force is attractive, i.e., spontaneous engulfment of reinforcing particles by the solidification front is expected, if ∆𝜎 < 0. Kaptay42 applied this criterion to a range of
𝑐𝑙𝑠
Al-ceramic systems (TiC, SiC, Al O , SiO and TiB ) and found that the Δσ values are
2 3 2 2 cls
positive for systems in which ceramic particles were observed to be pushed by the solidification front leading to particle segregation at the inter-grain area, which is undesirable for strength enhancement.
Due to the formation of an ordered C -like Mg layer on the exposed surfaces of NbC
NbC
particles, the effective interface between the liquid Mg and particle (Mg(l)/NbC) becomes Mg(l)/Mg at Mg@NbC. This implies a favoured wetting of the NbC
terminal
substrate by the molten Mg with a contact angle approaching zero (Θ =0). As a result,
clv
during solidification, the advancing Mg front interacts with the Mg@NbC particle to form a Mg(s)/Mg interface rather than a direct Mg(s)/NbC interface. As shown in Fig. 3a and Supplementary Fig. 13c, each terminal Mg atom exhibits 12 neighbours (four Mg atoms within the same plane, three Mg atoms in the 1st Mg layer, four Nb atoms and one C in the underlying NbC layer). Although crystallographically distinct, this face-centred-cubic-like coordination provides each Mg atom in the terminal layer with a coordination number equivalent to that of bulk Mg. Given this structural similarity, the surface energy of the Mg-coated NbC particle can be reasonably approximated as that of pure Mg. This leads to Δσ < 0, indicating spontaneous engulfment of NbC particles
cls
by the front during solidification of Mg(l)-Mg@NbC colloidal solution42. Furthermore, during solidification, Mg atoms at the advancing solidification front form Mg-Mg metallic bonds with the terminal Mg atoms on the surface of NbC particles, allowing the front to propagate through the Mg-coated NbC particles (Supplementary Fig.12). This behaviour indicates a weak correlation between the orientation of NbC particles and the solidification front, resulting in the engulfment of randomly oriented particles. Experimentally, no apparent orientation relationships between the NbC particles and the Mg matrix were observed (Supplementary Figs. 2, 4, 6).
It is also notable that NbC possesses a significantly higher shear modulus (~230 GPa43) than in-situ formed precipitates44-47. These engulfed, chemically bonded high modulus NbC particles within the Mg matrix directly contribute to the effective strengthening48 of the Mg matrix.
In conclusion, this combined atomistic-level experimental and theoretical study reveals the formation of an ordered Mg layer bonded to NbC particles (Mg@NbC), which is responsible for the observed spontaneous wetting and engulfment. The resulting material exhibits high specific tensile strength and stiffness, surpassing many high-performance metals. To identify other metal-ceramic systems where the base molten metal spontaneously wets the ceramic nanoparticles, future design should specifically employ ab initio interface structure prediction and interfacial energy calculations. We believe DFT calculation can help in accelerating ceramic particle selection for metallic materials by investigating their interfacial interaction. Beyond magnesium, these insights provide a blueprint for designing next-generation structural and functional metal-ceramic systems and scalable nanocolloid-based processing.

Methods
Fabrication of Mg-NbC materials. NbC powder feedstock with three particulate sizes were used in this study, and their morphology and size distribution are given in Supplementary Fig. 14. The NbC particles with the average size (D ) of 906 nm were
50
supplied from Companhia Brasileira de Metalurgia e Mineração (CBMM) coded as
NbC . The NbC particles with the average size (D ) of 287 nm were acquired from
micron 50
American Elements coded as NbC . The NbC particles with the average size (D )
submicron 50
of 15 nm were produced using plasma spray chemical synthesis49-51 coded as NbC
nano
(Supplementary Figs. 15, 16). To produce NbC
nano
, a thermal plasma jet was first generated with a hybrid plasma spray system using Ar and H gases. Raw Nb metal
2
powders with an average diameter of 45 µm (Kojundo Chemical Co. Ltd., Japan) were then introduced into the plasma to form high temperature Nb vapour through complete vaporisation. CH gas was also added to the plasma for carburisation of the Nb vapour
4
to form NbC particles. The feed rate of the raw powders and the CH gas flow rate were
4
adjusted to control the NbC nanoparticle size and Nb/C ratio of the NbC solid solution.
For particle size distribution measurements, using ImageJ 1.53k software52, ~1500
particles were analysed for each batch. The NbC and NbC powder
nano submicron
feedstocks were cold pressed into 16 mm diameter pellets using a Specac Atlas 15T
manual hydraulic press with 0.5 ton pressure. In the case of the NbC powder, 32
micron
mm diameter pellets were produced by applying 1 ton pressure. The pressed NbC
pellets were baked at 110°C for 12h and placed into a steel crucible containing molten
Mg (99.95% purity) at 680°C with a protective gas mixture of nitrogen and R134a gas in
the ratio of 12:1 (nitrogen flow at 3 L/min, R134a at 0.25 L/min), under ambient
pressure, to achieve pressure-less infiltration. During infiltration, the entrapped air in the
NbC precursor pellet escapes in the form of bubbles to the molten metal surface. The
infiltrated pellets were then solidified at 0.3 K/s cooling rate in protective gas. A pure Mg
reference sample was also solidified at the same cooling rate. The volume fraction of
the NbC particles in the Mg-NbC materials was calculated based on the measured
weight difference between the cold pressed NbC pellet and as-solidified Mg-NbC
materials, as well as the density of the pure Mg (1.78 g/cm3) and the NbC phase (7.8
g/cm3). The volume fraction of the NbC phase in as-solidified Mg-NbC , Mg-
nano
NbC , and Mg-NbC samples were 12.2 vol.%, 38.7 vol.%, and 53.8 vol.%,
submicron micron
respectively. A summary of the infiltrated materials is listed in Supplementary Table 1.
To demonstrate up-scalability for the preparation of a 1 kg batch Mg-NbC material, Mg-
38.7 vol.%NbC infiltrated pellets were preheated at 200°C for 2 h then added to a
submicron
Mg melt at 680°C under a protective atmosphere to obtain a final particle loading of 1
vol.%. After 15 min, the infiltrated Mg-NbC pellets were gently pressed into small
submicron
fragments and manually stirred to release and disperse the NbC particles into the
submicron
Mg melt. The resulting diluted Mg-1vol.%NbC colloid was poured into a steel
submicorn
mould and solidified at a cooling rate of approximately 30 K/s.
Mechanical characterisation. The micro-tensile specimens from as-solidified Mg-
NbC and Mg-NbC materials were prepared using a Thermo Fisher Scientific
nano submicron
Scios 2 DualBeam focused ion beam-scanning electron microscope (FIB-SEM) and
tested using a FemtoTools FT-NMT04 IN-SITU Nanoindentor under Field Emission-
SEM (FE-SEM) observation in the Scios 2 instrument (Supplementary Fig. 17). The
specimens were sectioned, lifted-out from the as-solidified Mg-NbC samples and then
attached to a tungsten needle by Pt deposition into a prefabricated hole whose size
conforms to the shape of the lifted sample. This was the fixed side in the tensile testing
rig. The specimen was shaped into a dog-bone geometry with rectangular cross section
(~10 m gauge length and 12 m2 cross sectional area). The dog-bone specimen was
attached to a Si gripper that connected to a microforce sensing probe (FT-S200’000)
within FT-NNMT04. The micro-tensile test was performed in the displacement control
mode and a strain rate of 110-4 s-1. The video recording of the in-situ tensile tests is
presented in supplementary files (Supplementary Videos 1-3).
The macro-tensile tests of the as-solidified Mg-1vol.% NbC and Mg-
submicron
13vol.%NbC materials and reference pure Mg were conducted on a Instron 5500
micron
Universal Electromechanical Testing System with a constant cross head speed of 0.225
mm/min (2×10−4 s−1 initial strain rate) and the results are shown in Supplementary Fig.
7b and Supplementary Note Fig. 2c.
A Micro Materials NanoTest ALPHA with a Berkovich tip was used for nanoindentation
tests with load-control to a maximum load of 100 mN with a 6.7 mN/s loading rate, held
for 10 s, and unloaded at 6.7 mN/s. Hardness and elastic modulus are evaluated from
the load-indentation depth curves.
Strengthening contribution calculations. The two synergetic strengthening
mechanisms in this work are particle-dislocation interaction (Orowan strengthening) and
load-transfer strengthening. The Orowan strengthening13 can be expressed as:
𝜑𝐺 𝑏 6𝑉 1/3
𝑚 𝑃
∆𝜎 = ( )
𝑂𝑟𝑜𝑤𝑎𝑛 𝑑 𝑝 𝜋 S (3)
Where 𝜑 is a coefficient with a value of 2. 𝐺 is the shear modulus of the matrix and is
𝑚
17.3 GPa for Mg53. b is the Burgers vector of Mg and has value of 0.32 nm54. 𝑉 denotes
𝑝
the volume fraction of NbC particles, and 𝑑 represents the NbC particle size. For each
𝑝
NbC catalogue, the d , d , d particle sizes (Supplementary Table 1) were used
peak- peak peak+
to calculate the Orowan strengthening contribution.
The load-transfer strengthening increment is given by,
𝛥𝜎 = ( 𝐸𝑝 −1)𝑉 𝜎𝑚 (4)
𝐿𝑇 𝐸𝑚 𝑝 𝑌𝑆
Where 𝐸 is the elastic modulus of NbC and is taken as 480 GPa43 and 𝐸 is the elastic
𝑃 𝑚
modulus of the Mg matrix and is taken as 45 GPa53. 𝜎𝑚 is the measured yield strength
𝑌𝑆
of Mg (35 MPa). 𝑉 is the volume fractions of the NbC particles. Further details of load-
𝑝
transfer contribution, 𝛥𝜎 is given in Supplementary Note 2.
𝐿𝑇,
Grain boundary strengthening is considered for the polycrystalline Mg-NbC materials
evaluated using macro-tensile test. The yield strength improvement arising from the
−1/2 −1/2
Hall-Petch relationship55, 56 can be written as ∆𝜎 = 𝑘∙𝑑 −𝑘 ∙𝑑 , where, 𝑘 is
𝐺𝑅 𝑀𝑔−𝑁𝑏𝐶 𝑀𝑔
the Hall-Petch coefficient57 with value of 150 MPa·m-1/2, 𝑑 and 𝑑 are the
𝑀𝑔−𝑁𝑏𝐶 𝑀𝑔
average Mg grain size with and without NbC addition. The resulting ∆𝜎 is
𝐺𝑅
approximately 4.9 MPa for Mg-1 vol.%NbC material (𝑑 = 149 m and 𝑑
submicron 𝑀𝑔−𝑁𝑏𝐶 𝑀𝑔
= 417 m), and approximately 2.6 MPa for Mg-13 vol.%NbC material (𝑑 is
micron 𝑀𝑔−𝑁𝑏𝐶
1.1 mm and 𝑑 is 5.0 mm).
𝑀𝑔
Microstructure observations. A Zeiss Crossbeam 340 SEM and Thermo Fisher
Scientific Talos F200X were used to investigate the microstructure of the NbC powder
feedstock and the Mg-NbC samples. Scanning transmission electron microscopy
(STEM) images and energy dispersive x-ray spectroscopy (EDS) elemental maps were
collected using the Talos F200X instrument.
Atomic-resolution STEM imaging and electron energy loss spectroscopy (EELS)
mapping were conducted using an aberration corrected Nion UltraSTEM100 working at
100 kV. The microscope is equipped with a cold-FEG emitter with 0.3 eV native energy
spread. With a fifth-order probe aberration corrector, the microscope optics can be
configured to enable a probe size of < 1 Å at 100 kV with a convergence semi-angle of
30 mrad and a probe current of 30 pA in the conditions used for these experiments. The
semi-angular ranges for high-angle annular-dark-field (HAADF), medium-angle annular-
dark-field (MAADF) images and annular-bright-field (ABF) images were 89 -195 mrad,
52-89 mrad (or 30-195 mrad when acquired simultaneously with ABF images), and 10-
30 mrad, respectively. Electron energy loss spectra were acquired on a Gatan Enfina
spectrometer that is retrofitted with a Quantum Detectors Merlin EELS hybrid pixel
camera. EELS spectrum images (SIs) were acquired in “event-streamed” mode using a
36 mrad collection semi-angle. Leveraging the instrument’s high stability, characterised
by negligible sample drift under experimental conditions, multiple consecutive SI scans
with short pixel dwell times (3 or 5 ms/pixel) were accumulated and averaged. This
approach enabled sufficient signal acquisition while minimizing noise and sample
damage through repeated acquisitions. Prior to data processing, the EELS SI dataset
was denoised via the built-in principal component analysis (PCA) function in Gatan
Microscopy Suite 3.6 (GMS 3.6), with careful inspection of residuals to ensure artifact-
free reconstruction. Notably, the near-ideal Poisson noise characteristics of data
collected using next generation hybrid pixel detectors make them particularly well-suited
for PCA-based denoising, enabling enhanced signal extraction with minimal artifact
introduction58. EELS elemental maps were generated in two ways. For SIs containing
the Nb M edges (205 eV) and C K (284 eV) in close proximity, we used model-based
EELS quantification function in GMS 3.6 to separate the two signal components for
elemental mapping. This method was applied to create Fig. 1b, Supplementary Fig. 15d
and Fig. 16d, and the C K map in Fig. 3c. For the Mg K (1305 eV, in Fig. 3c) and Nb
L (2465 eV, in Fig. 3c, Supplementary Fig. 15e and Fig. 16e) edges, we first
2,3
subtracted a decaying background preceding each edge using a fitted power-law
function and then integrated the EELS edge intensity over a 50 eV window from each
edge onset. EELS map intensities are shown using a false colour scale, where low
intensity indicates lower relative elemental concentration, and higher contrast reflects a
relative increase in concentration. We note that two consecutive SIs were acquired in
the same region in Fig. 3c to cover a wide energy range involving all the
aforementioned core-loss edges, while only C K, Nb L and Mg K were selected for
2,3
illustration in the main text.
Cross-sections of the bulk samples were prepared using standard metallographic
methods with the oil-based diamond suspension. TEM specimens from the Mg-NbC
samples were initially sliced from the bulk and ground to a thickness of ~100 µm,
followed by dimpling using a Gatan dimple grinder system and subsequent thinning
using ion milling with a Gatan PIPS 695 instrument that operated at 3-5 kV and beam
angles of 3-5°. To remove damaged layers and any amorphous structure on the TEM
foil surface, a final polishing step was performed at 1 kV with Ar⁺ ions at a 7° incident
beam angle. Electron backscattered diffraction (EBSD) was performed to characterise
the grain size of Mg in Mg-NbC samples. Specimens were placed in the Tescan Magna
UHR-SEM tilted 70° from the horizontal plane towards the EBSD camera with a 30 kV
electron beam, 13 nm spot size and 80-100 nm step size. The 4D-STEM based
orientation/phase mapping was applied to produce crystallographic maps of the Mg-
NbC specimen at high resolution using a Talos F200X and NanoMegas ASTAR
nano
system at 200 kV, 2 nm step size. The data analysis was conducted by ATEX 5.0
software (Academic Edition)59.
The X-ray diffraction (XRD) spectrum for Mg-13vol.%NbC sample was acquired
micron
using a Bruker D8 ADVANCE XRD diffractometer with a Cu K (= 0.1542 nm) radiation

source, and the results are shown in Supplementary Note Fig. 2b. The candidate
phases for matching are selected from Powder Diffraction File™ (PDF®) database.
DFT calculations. The present ab initio molecular dynamics (AIMD) study employs the
first-principles code VASP (Vienna Ab initio Simulation Package)60 that utilizes the
periodic boundary conditions (PBC). Supercells were built for modelling the interfaces
between liquid Mg and NbC substrates. NbC {001} substrate has an equal number of
Nb and C atoms at the surface, being electronically neutral and non-polar. The number
of liquid Mg atoms for the systems is fixed to be 300, while the number of Nb atoms and
C atoms are fixed to be 72. A tetragonal supercell was created with a = 3a for the
0
Mg(l)/NbC interface, where a is the length of the axis of the conventional NbC cell
{001} 0
with consideration of the thermal expansion at the simulation temperature61. The length
of the c-axis is determined by the thickness of the NbC slab and the volume of the Mg
atoms with the density at the simulation temperature61. The built supercells are
summarized in Supplementary Table 4 and the related configurations are shown in
Supplementary Fig. 18a. To avoid artificial interface interactions, a space of about 2.0 Å
was made between the substrate atoms and the liquid Mg atoms.
VASP employs the electronic density-functional theory (DFT) within the projector-
augmented wave framework62. CThe generalised gradient approximation (GGA-PBE) was
used for the exchange and correlation terms63. The cut-off energy for the ab initio
molecular dynamics simulations is 320.0 eV and Г-point in the Brillouin zone is used as
there is a lack of symmetry in the liquid/solid interfaces60, 64. The simulation temperature
is set to be 1000 K.
A two-step approach was used: first AIMD simulations were performed with the
substrate atoms fixed for about 300 steps (1.5 femtosecond (fs) per step). Then, the
AIMD simulations are continued with relaxation of all the atoms. The simulations
revealed thermal equilibration after about 1 picosecond (ps) (see Supplementary Fig.
18). This approach avoids risks of collective movements of atoms during the AIMD
simulations31.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "The manuscript titled “Ultrahigh strength magnesium via solidification of nanocolloid” presents a comprehensive study on\nMg–NbC composites produced using NbC particles of varying sizes (nano, submicron, and micron). The work combines\nsolidification experiments, detailed electron microscopy, and ab initio molecular dynamics simulations to investigate\ncolloidal stability and interfacial bonding. The technical depth and breadth of data presented, particularly in the\nSupplementary Information, are impressive.\nHowever, some critical issues need clarification:\n1. The manuscript repeatedly compares the strengthenability of composites reinforced with different NbC particle sizes, but\nthis term is not rigorously defined. The key experimental comparison should focus on how the absolute mechanical\nproperties—yield strength, ductility, and modulus—evolve with particle size and volume fraction. If “strengthenability” refers\nto strengthening efficiency, the authors must clearly define the normalization basis (e.g., per particle surface area, interface\nvolume, or particle-matrix load transfer coefficient).\n2. Given that the authors invoke Mg@NbC coating and CNbC-like interfacial layers, the particle-size dependence should\nnaturally be linked to surface-to-volume ratio and interfacial energetics. Yet, no direct evidence or quantitative discussion of\nhow surface energy or bonding characteristics evolve with particle size is provided. A clear correlation between AIMD-\nderived interface stability and experimental mechanical performance is missing.\n3. Table S2 shows that the NbC volume fraction increases dramatically with particle size—from 12% (nano) to ~54%\n(micron). This large variation introduces a confounding factor: it is unclear whether the observed strengthening trend arises\nprimarily from particle size, particle fraction, or both. A normalization or controlled comparison (e.g., constant volume fraction)\nwould help isolate the dominant strengthening mechanism.\n4. The manuscript attributes strength enhancement to size effects and interface ordering, yet contributions from grain\nrefinement, load transfer, and dislocation–particle interactions are not separated. Clarification through composite modeling\nor a rule-of-mixtures comparison would enhance the credibility of the claims.\n5. The paper is data-rich but difficult to navigate. The key mechanistic insights—particularly the transition from colloidal\nstability to interface-controlled strengthening—should be summarized more clearly within the main text rather than being\ndeferred to supplementary sections.\nOverall, this is a technically strong and potentially impactful study, but the mechanistic rationale linking particle size,\ninterface stability, and mechanical strengthening can be improved.", "reviewer_2": "This paper presents an innovative approach to producing ultrahigh-strength magnesium by infiltrating molten Mg into a pellet\nof NbC nanoparticles. It combined atomic-scale experimental evidence and first-principles DFT simulations to demonstrate\nthe spontaneous wetting and coherent bonding at the Mg/NbC interface. The fabrication method is fairly straight forward. A\nuniform dispersion of NbC nanoparticles without relying on external fields or complex thermomechanical processing was\nachieved.\nHowever, the powder pellet method is somewhat difficult to scale up for large scale production. The pellet method may limit\nits flexibility of NbC particle volume percent. Long-term thermal stability of the Mg-NbC composites will be of interest.\nOverall, it is a plausible paper with merit.", "reviewer_3": "The authors have undertaken a very interesting study on the role of dispersed nano-particles on the properties of Mg, which\nthey are calling a nano-colloid. This has been the challenge with nano-composites, i.e. how to achieve a homogeneous\ndispersion of nano-particles, rather than finding them in clusters. I am aware of such processes having been used previously\nto create nano-composites. Some of the challenge can be when these pellets are added to melts particularly with gravity\nseparation because normally lower levels of particles are preferred.\nOne of the drawbacks of this paper is that it compares the properties of these materials with standard alloys. However, what\nis shown here are metal matrix compositions with a very large amount of the second phase. The volume % are 12, 39 and\n54%, however it should be realised that given the density different the mass% is a lot greater. This is the primary reason for\nthe increase in modulus and hardness, i.e. that is it not really a Mg based material, it is primarily a ceramic. Having said that\nthe nano-particle version still is majority Mg although it is still high enough content to be a metal matrix composite and\nshould be compared with other MMCs. So comparison with similar materials is required.\nI also do not believe that the micro-tensile tests are providing a useful comparison of the properties. Samples at this size are\nwell below critical flaw sizes and consequently do not give a reasonable estimate of the tensile properties that can be\nexpected in typical use cases. Nor should these results be compared with results from macro-tensile tests. The failure is\ncompletely different as shown by the tensile curve results. The failure is seen in Supp Fig 7 which is interesting but not\nhelpful for comparision with bulk properties. The real tensile data is observed in Supplementary Fig 2 which is very different.\nSo comparison of materials in the same or similar test is required.\nI think that when appropriate materials and testing methods are compared the properties shown are not as outstanding as it\nappears.\nThere is a nice microscopic and DFT study of the interface indicating that there is templating on the particles as has been\ndescribed previously by researchers from Brunel in this group. Whilst DFT is a useful tool, it is not the only tool to\nunderstanding the interface behaviour. The conclusions indicate that DFT can predict the wetting behaviour of metals on\nceramic nano-particles. It is not clear to me that NbC was chosen because of the these calculations. If not, I would be\ninterested in seeing results that backed that statement up more strongly.\nLine 246 terminal is misspelled.\nSupp Fig 8 has misspelled label on y-axis."}, "author_response": "Authors response to reviewers’ comments (NCOMMS-25-78483-T)\nReviewer #1 (Remarks to the Author):\nThe manuscript titled “Ultrahigh strength magnesium via solidification of nanocolloid”\npresents a comprehensive study on Mg–NbC composites produced using NbC particles of\nvarying sizes (nano, submicron, and micron). The work combines solidification experiments,\ndetailed electron microscopy, and ab initio molecular dynamics simulations to investigate\ncolloidal stability and interfacial bonding. The technical depth and breadth of data presented,\nparticularly in the Supplementary Information, are impressive.\nOverall, this is a technically strong and potentially impactful study, but the mechanistic\nrationale linking particle size, interface stability, and mechanical strengthening can be\nimproved.\nWe thank reviewer for identifying the strengths and seeking further clarification on the\nfollowing points:\nR1.1. The manuscript repeatedly compares the strengthenability of composites reinforced\nwith different NbC particle sizes, but this term is not rigorously defined. The key experimental\ncomparison should focus on how the absolute mechanical properties—yield strength,\nductility, and modulus—evolve with particle size and volume fraction. If “strengthenability”\nrefers to strengthening efficiency, the authors must clearly define the normalization basis\n(e.g., per particle surface area, interface volume, or particle-matrix load transfer coefficient).\nResponse: We thank the reviewer for highlighting the ambiguity associated with the term\n“strengthening effectiveness (strengthenability)” in the original manuscript. We agree that,\nwithout a rigorous definition and clear normalisation basis, this terminology could lead to\nconfusion when comparing composites reinforced with different particle sizes and volume\nfractions.\nIn the revised manuscript, we have therefore removed the use of the term “strengthening\neffectiveness (strengthenability)”. Instead, the discussion now focuses explicitly on the\nabsolute evolution of mechanical properties, particularly the yield-strength increment, Δσ\nYS\nas a function of NbC particle size and volume fraction (Fig. 2b). This approach provides a\ndirect and unambiguous basis for comparison, consistent with the reviewer’s\nrecommendation.\nTo clearly separate the effects of particle size and volume fraction on yield strength\nenhancement, we have added a new plot (Fig. 2b) showing Δσ as a function of NbC\nYS\nvolume fraction for all particle sizes investigated. In addition, the individual strengthening\ncontributions arising from particle-dislocation interactions (Orowan strengthening) and load\ntransfer are explicitly evaluated using particle size dependent parameters. The governing\nexpressions for these mechanisms are detailed in the Methods section, with an expanded\ndiscussion of the load-transfer mechanism provided in Supplementary Note 2.\nBy presenting both the experimentally measured yield strength increments and the\ncorresponding mechanistic contributions as functions of particle size and volume fraction,\nthe revised manuscript provides a physically meaningful and consistent framework for\ncomparing strengthening behaviour, without reliance on ambiguous terminology. Accordingly,\nthe statement in the original manuscript reporting specific strengthening effectiveness in unit\nof MPa/vol% for different particle sizes has been removed.\nR1.2. Given that the authors invoke Mg@NbC coating and CNbC-like interfacial layers, the\nparticle-size dependence should naturally be linked to surface-to-volume ratio and interfacial\nenergetics. Yet, no direct evidence or quantitative discussion of how surface energy or\nbonding characteristics evolve with particle size is provided. A clear correlation between\nAIMD-derived interface stability and experimental mechanical performance is missing.\nResponse: We agree with the reviewer that the particle-size dependence of strengthening\nshould be quantitatively linked to surface-to-volume effects and interfacial energetics and\nfurther connected to a physically meaningful strengthening framework. In the revised\nmanuscript, we now explicitly connect AIMD-derived Mg-C bonding, particle-size-dependent\ninterfacial atom fraction, interface-limited debonding strength, and load-transfer\nstrengthening to the experimentally observed mechanical response (Supplementary Note 2).\nAb initio molecular dynamics simulations reveal strong chemical bonding between superficial\nC atoms on NbC and terminating Mg atoms, forming ordered Mg-C interfacial bonds (Fig. 4).\nThese predictions are consistent with STEM observations of an interfacial multilayer\ncomplexion that is epitaxially coherent with the NbC surface (Fig. 3), indicating the formation\nof Mg-coated NbC particles (Mg@NbC) in the liquid Mg that are retained upon solidification.\nThe AIMD calculations further yield an interfacial fracture energy of Γ = 1.54 J·m-2, providing\na quantitative metric of interface stability.\nTo translate this interfacial energetics into a mechanical limit for stress transfer, we evaluated\nthe interface-limited debonding strength using a fracture-energy-based criterion for a stiff\nspherical inclusion embedded in a compliant matrixs6:\n1 16 Γ 𝐸 (1+𝜈 )\n𝜎 (𝑑) = √ 𝑚 𝑚\ndebond 3(1−𝜈 ) 𝑑\n𝑚\nwhere 𝐸 and 𝜈 are the Young’s modulus and Poisson’s ratio of Mg. Using 𝐸 = 45GPa,\n𝑚 𝑚 𝑚\n𝜈 = 0.35, and Γ = 1.54 J·m-2, we obtain predicted debonding stresses of approximately\n𝑚\n5.12 GPa (15 nm), 1.17 GPa (287 nm), and 0.66 GPa (906 nm). These gigapascal-level\nvalues indicate that the Mg-NbC interface is sufficiently strong to sustain stress transfer\nacross all particle sizes studied, with increasing resistance to debonding as particle size\ndecreases.\nIn parallel, the particle-size dependence of the population of chemically bonded interfacial\natoms was estimated by considering the ratio of surface to bulk C atoms in a cubic NbC\nparticle with {001} surfaces:\n𝑛(𝐶 ) 6𝑑2/𝑎2 1.5\nsurf\n𝑅 = = =\n𝑛(𝐶 ) 4𝑑3/𝑎3 𝑑/𝑎\nbulk\nwhere 𝑎 = 4.469 Å is the NbC lattice parameter and 𝑑 is the particle size. This scaling shows\nthat the fraction of surface atoms decreases rapidly with increasing particle size, providing a\nquantitative explanation for why interfacial chemistry is expected to contribute most strongly\nat the nanoscale.\nFinally, we connect these interfacial considerations to load transfer strengthening. Under the\niso-strain (Voigt) load-sharing condition appropriate for well-bonded particulate composites,\nthe load transfer strengthening contributions4, s5 is:\n𝐸\nΔ𝜎 = ( 𝑝 −1)𝑉 𝜎𝑚\nLT,Voigt 𝐸 𝑝 𝑌𝑆\n𝑚\nwhere 𝐸 is the NbC modulus, 𝑉 is the particle volume fraction, and 𝜎𝑚 is the yield strength\n𝑝 𝑝 𝑌𝑆\nof the unreinforced Mg matrix processed under identical conditions (𝜎𝑚 = 35 MPa).\n𝑌𝑆\nUsing representative moduli 𝐸 /𝐸 ≈ 10, this reduces to Δ𝜎 = 9𝑉 𝜎𝑚. For the three\n𝑝 𝑚 LT,Voigt 𝑝 𝑌𝑆\ncompositions corresponding to the particle sizes studied here, the predicted load transfer\ncontributions are:\n• 15 nm, 12.2 vol.% NbC: Δ𝜎 = 9(0.122)(35) ≈ 38.4MPa\nLT\n• 287 nm, 38.7 vol.% NbC: Δ𝜎 = 9(0.387)(35) ≈ 121.9MPa\nLT\n• 906 nm, 13 vol.% NbC: Δ𝜎 = 9(0.13)(35) ≈ 41.0MPa\nLT\nCrucially, the interface must be able to sustain these stress-transfer demands. A\nconservative interface-capacity estimate may be written as Δ𝜎 ∼ 𝜂𝑉 𝜎 (𝑑)with 𝜂 ≈\nLT,max 𝑝 debond\n0.5. Using the above 𝜎 values, the interface-limited capacities remain far above the\ndebond\nVoigt load transfer requirements for all three cases, confirming that the Mg-NbC system\noperates in the perfect load transfer (Voigt-limited) regime. Therefore, for the present Mg-\nNbC materials, particle size does not limit load transfer through debonding; instead, particle\nsize primarily modulates the relative volumetric significance of interfacial chemistry (via the\ninterfacial atom fraction 𝑅), while the overall macroscopic strengthening is governed by the\ncombined effects of particle fraction (load transfer) and dislocation-mediated Orowan\nmechanisms.\nThe detailed load transfer increment (Δ𝜎 ) assessment considering both the interface-\nLT\nlimited debonding and Voigt load-sharing mechanisms are presented in Supplementary Note\n2.\nReferences\ns4. Taya, M. & Arsenault, R. J. in Metal matrix composites: thermomechanical behavior\n(Elsevier, 2016).\ns5. Withers, P. J. & Clarke, A. P. A neutron diffraction study of load partitioning in continuous\nTi/SiC composites. Acta Materialia 46, 6585–6598 (1998).\ns6. Park, K. & Paulino, G. H. Cohesive zone models: a critical review of traction-separation\nrelationships across fracture surfaces. Appl. Mech. Rev. 64, 060802 (2011).\nR1.3. Table S2 shows that the NbC volume fraction increases dramatically with particle\nsize—from 12% (nano) to ~54% (micron). This large variation introduces a confounding\nfactor: it is unclear whether the observed strengthening trend arises primarily from particle\nsize, particle fraction, or both. A normalization or controlled comparison (e.g., constant\nvolume fraction) would help isolate the dominant strengthening mechanism.\nResponse: We agree with the reviewer that the effect of the particle size and particle\nfraction should be explicitly compared to understand the participating strengthening\nmechanisms. To address this, we have added Fig. 2b, which presents the strengthening\ncontribution as a function of NbC volume fraction for all three particle sizes studied. This\nnormalisation allows the effect of particle size to be separated from that of particle fraction.\nFor example, at a volume fraction of approximately 13%, the calculations indicate that\nsmaller particle sizes correspond to higher yield-strength increments. This trend is\nexemplified by comparing the Mg-NbC (12.2 vol.%) and Mg-NbC (13 vol.%) samples,\nnano micron\nwhich exhibit similar predicted load-transfer contributions (see the response to R1.2 above)\nbut markedly different Orowan contributions due to the particle-size effect (see the response\nto R1.4 below).\nFurthermore, in the inset to Fig. 2b, we now compare the yield strength improvement of the\npresent Mg-NbC materials with a wide range of metal matrix composites produced by melt\nprocessing, powder metallurgy, and additive manufacturing routes. This contextual\ncomparison highlights the exceptional strengthening contribution achieved via the colloidal\nsolidification approach.\nR1.4. The manuscript attributes strength enhancement to size effects and interface ordering,\nyet contributions from grain refinement, load transfer, and dislocation–particle interactions\nare not separated. Clarification through composite modelling or a rule-of-mixtures\ncomparison would enhance the credibility of the claims.\nResponse: Grain refinement can be excluded for the micro-tensile test as the length scale of\nsample is ~10 m, while the corresponding Mg grain size of the same sample is > 150 m.\nThe dominant strengthening mechanisms in this work are therefore particle-dislocation\ninteractions (Orowan strengthening) and load transfer. The Orowan contribution is now\nexplicitly calculated using established models and is presented in Fig. R1 as a function of\nNbC volume fraction for all particle sizes.\nThe load-transfer contribution, originating from stress transfer across the Mg-NbC interface,\nis calculated using a composite mechanics framework and is presented in Fig. R2. By\ncombining the contributions shown in Figs. R1 and R2, the theoretical prediction of the yield\nstrength improvement of the Mg-NbC system is obtained and presented in Fig. 2b in the\nrevised manuscript.\nA section named “Strengthening contribution calculations” section is now added to the\nMethods in the revised manuscript. It includes the details of the Orowan strengthening and\nload transfer strengthening calculations.\nFigure R1. Calculated yield strength improvement (∆𝜎 ) of Mg-NbC materials via Orowan\n𝑌𝑆\nstrengthening with nano, submicron, and micron sized particles. The dotted, solid and dashed lines\ncorrespond to the calculation with d d , and d particle size.\npeak- peak peak+\nFigure R2. Calculated yield strength improvement (∆𝜎 ) of Mg-NbC materials via Load transfer\n𝑌𝑆\nstrengthening as a function of the particle volume fraction.\nR1.5. The paper is data-rich but difficult to navigate. The key mechanistic insights—\nparticularly the transition from colloidal stability to interface-controlled strengthening—should\nbe summarized more clearly within the main text rather than being deferred to\nsupplementary sections.\nResponse: We have reorganised the manuscript accordingly. The discussion of colloidal\nstability, which is critical for achieving uniform dispersion during solidification, has been\nmoved from the Supplementary Information to the main text in the revised manuscript. This\nis seen in the section “Stable Mg-NbC nanocolloid at high temperature” on page 13.\nReviewer #2 (Remarks to the Author):\nThis paper presents an innovative approach to producing ultrahigh-strength magnesium by\ninfiltrating molten Mg into a pellet of NbC nanoparticles. It combined atomic-scale\nexperimental evidence and first-principles DFT simulations to demonstrate the spontaneous\nwetting and coherent bonding at the Mg/NbC interface. The fabrication method is fairly\nstraight forward. A uniform dispersion of NbC nanoparticles without relying on external fields\nor complex thermomechanical processing was achieved.\nOverall, it is a plausible paper with merit\nWe thank the reviewer for their positive assessment of the manuscript and for recognising\nthe novelty of the processing strategy, the atomic-scale experimental evidence, and the\nintegration of first-principles simulations. We also appreciate the reviewer’s constructive\ncomments regarding scalability and thermal stability, which we address below:\nR2.1. The powder pellet method is somewhat difficult to scale up for large scale production.\nThe pellet method may limit its flexibility of NbC particle volume percent.\nResponse: We appreciate this important point. The original motivation for developing the\npellet infiltration process was to enable the safe and controlled introduction of nanoparticles\ninto molten Mg, particularly in industrial environments where direct handling of nano powders\nis challenging. Importantly, the infiltration process (Fig. 1) produces a highly concentrated\nMg-NbC master alloy, rather than a final-composition material.\nThis master alloy can subsequently be diluted to any desired particle volume fraction using\nconventional melt processing, making the approach compatible with standard foundry\npractice. In the revised manuscript, we now explicitly demonstrate this concept by reporting\nthe dilution of a concentrated Mg-NbC master alloy to produce bulk material with 1\nsubmicron\nvol.% NbC particle loading. Corresponding macroscopic billet, tensile specimens, and\nstress-strain curves are now included (Supplementary Fig. 7), demonstrating both scalability\nand compositional flexibility of the process.\nR2.2. Long-term thermal stability of the Mg-NbC composites will be of interest.\nResponse: We agree with the reviewer that long-term thermal stability is an important\nconsideration. One anticipated advantage of the present Mg-NbC system is the use of a\nthermally stable ceramic phase, in contrast to conventional Mg alloys where strengthening\nrelies on temperature-sensitive precipitates. While a comprehensive high-temperature\nmechanical and microstructural stability study is beyond the scope of the present\nmanuscript, such investigations are currently ongoing and will be reported in a future\ndedicated publication once completed. This aspect has been included in this revised\nmanuscript.\nReviewer #3 (Remarks to the Author):\nThe authors have undertaken a very interesting study on the role of dispersed nano-particles\non the properties of Mg, which they are calling a nano-colloid. This has been the challenge\nwith nano-composites, i.e. how to achieve a homogeneous dispersion of nano-particles,\nrather than finding them in clusters. I am aware of such processes having been used\npreviously to create nano-composites. Some of the challenge can be when these pellets\nare added to melts particularly with gravity separation because normally lower levels of particles\nare preferred.\nWe thank the reviewer for the thoughtful and constructive assessment of our work, and for\nrecognising the long-standing challenge of achieving homogeneous dispersion of\nnanoparticles in metal matrix composites. We appreciate the opportunity to clarify the scope,\ncomparisons, and interpretation of our results.\nIn the present work, in addition high volume particle loading, we also demonstrate that\nhomogeneous dispersion at lower particle loading (~1 vol.% NbC ) can be achieved\nsubmicron\nusing an infiltrated Mg-NbC master pellet (Supplementary Fig.7), which is introduced into the\nmelt under controlled processing conditions. This approach mitigates particle clustering by\npreventing direct particle-particle contact during addition and promotes uniform distribution\nthrough wetting-assisted dispersion and melt convection. Furthermore, the discussion in\n“Assessment of colloid stability for larger particles” in Supplementary Note 1 addresses the\nstability of Mg-NbC colloid in the presence of gravity.\nR3.1. One of the drawbacks of this paper is that it compares the properties of these\nmaterials with standard alloys. However, what is shown here are metal matrix compositions\nwith a very large amount of the second phase. The volume % are 12, 39 and 54%, however\nit should be realised that given the density different the mass% is a lot greater. This is the\nprimary reason for the increase in modulus and hardness, i.e. that is it not really a Mg based\nmaterial, it is primarily a ceramic. Having said that the nano-particle version still is majority\nMg although it is still high enough content to be a metal matrix composite and should be\ncompared with other MMCs. So comparison with similar materials is required.\nResponse: We agree with the reviewer that materials with high reinforcement volume\nfractions should be compared primarily with other metal matrix composites. To address this\npoint, we have expanded the comparative framework in the revised manuscript.\nSpecifically, we have collated data from the literature on Mg-, Al-, Ti-, and Fe-based metal\nmatrix composites reinforced with oxides, carbides, borides, and nitrides, which is now\npresented in Supplementary Tables 1 and 3. A direct comparison between these established\nMMC systems and the present Mg-NbC materials is provided as inset in Fig. 2b. This\ncomparison shows that, even when benchmarked against other MMCs with comparable\nreinforcement levels, the Mg-NbC system processed using simple solidification process\nexhibits remarkable improvement in yield strength. The dataset in Fig. 2d compares strength\nand modulus relative to density of a range of alloys and their composites.\nR3.2. I also do not believe that the micro-tensile tests are providing a useful comparison of\nthe properties. Samples at this size are well below critical flaw sizes and consequently do\nnot give a reasonable estimate of the tensile properties that can be expected in typical use\ncases. Nor should these results be compared with results from macro-tensile tests. The\nfailure is completely different as shown by the tensile curve results. The failure is seen in\nSupp Fig 7 which is interesting but not helpful for comparison with bulk properties. The real\ntensile data is observed in Supplementary Fig 2 which is very different. So comparison of\nmaterials in the same or similar test is required.\nI think that when appropriate materials and testing methods are compared the properties\nshown are not as outstanding as it appears.\nResponse: We agree that micro-tensile and macro-tensile tests probe different aspects of\nmechanical behaviour and should be interpreted accordingly. We have clarified this\ndistinction more explicitly in the revised manuscript. In the section “Ultrahigh tensile strength\nin Mg-NbC materials” on page 6, the Fig.2 has been modified and expanded by adding new\nresults, and the corresponding text is also revised.\nIn this work, micro-tensile testing (Fig. 2a) was intentionally performed on materials with a\nsingle-crystal Mg matrix, both without particles and with dispersed NbC particles. This design\nallows a direct and controlled comparison of the intrinsic strengthening effect of NbC\nparticles on Mg deformation, while eliminating contributions from grain boundaries and, texture. In this context, the micro-tensile results are not intended to represent bulk\nengineering properties, but rather to reveal the fundamental strengthening capability of the\nMg-NbC system at the crystal level.\nIn contrast, newly added Fig. 2b includes two distinct datasets: (i) micro-tensile results\nobtained from Mg-NbC material (square symbol), and (ii) macro-tensile results obtained from\npolycrystalline Mg-NbC materials (diamond symbol) containing submicron and micron-sized\nNbC particles, following ASTM E8 standard. These macro-tensile results reflect bulk\nmechanical behaviour and incorporate the effects of grain structure, particle distribution, and\nprocessing-induced defects.\nThe two testing approaches are used in a complementary manner (micro-tensile tests\nestablish the intrinsic, upper-bound strengthening potential arising from particle-matrix\ninteractions, while macro-tensile tests demonstrate the translation of this potential to bulk,\npolycrystalline materials at realistic length scales). The yield strength improvement\nmeasured in macro and micro-tensile experiments are consistent with the theoretical\npredicted tendency for all particle size and volume fraction. Dataset in Fig. 2b, provide a\ncoherent picture of mechanical response at both local and bulk Mg-NbC materials.\nR3.3. There is a nice microscopic and DFT study of the interface indicating that there is\ntemplating on the particles as has been described previously by researchers from Brunel in\nthis group. Whilst DFT is a useful tool, it is not the only tool to understanding the interface\nbehaviour. The conclusions indicate that DFT can predict the wetting behaviour of metals on\nceramic nano-particles. It is not clear to me that NbC was chosen because of the these\ncalculations. If not, I would be interested in seeing results that backed that statement up\nmore strongly.\nResponse: We first experimentally observed the exceptional wetting characteristic and\ncolloidal stability of NbC in liquid Mg. The ab initio molecular dynamics simulations were\nsubsequently performed to rationalise and interpret this experimentally observed\nphenomenon at the atomic scale. The simulations reveal the formation of ordered Mg-C\ninterfacial bonding and provide a quantitative measure of interfacial stability, supporting the\nexperimental observations. We believe DFT calculation can help in accelerating ceramic\nphase selection for metallic materials by investigating the formation of ordered layer on\nceramic phase as one of the necessary conditions to realise the excellent wetting of the\nmolten metal to the ex-situ incorporated ceramic phases.\nR3.4. Line 246 terminal is misspelled.\nSupp Fig 8 has misspelled label on y-axis.\nResponse: All identified typographical issues have been corrected, and a comprehensive\nproofreading of the Manuscript and Supplementary Information has been conducted to\nimprove overall clarity and presentation."}, {"round": 2, "reviewer_comments": {"reviewer_1": "The authors have thoroughly addressed my previous comments, resulting in substantial improvements to the manuscript.\nThe work is now complete and suitable for publication in its current form.", "reviewer_2": "It seems the authors addressed my comments reasonably well. The dilution of master nanocomposites into magnesium\nalloys shows good results too.", "reviewer_3": "I still have a lot of concerns about how the results are presented, particularly with the use of micro-tensile testing. The values\nshould not be compared with other sanples where macro-testing is used. This gives a false impression that the properties\nare a lot better than they actually are. In SSupplementary Figure 2 we see the Mg14%NbC alloy has a yield strength of\napproximately 150MPa with an elongation to fracture of about 1%. This is typical for alloys such as this one. This is very\ndifferent from the 678MPa yield strength improvement mentioned in the manuscript. If this is only compared with other micro\ntensile tests that would be fine, but that does not seem to be the case."}, "author_response": "Response to Reviewer #3’s Comments (NCOMMS-25-78483A)\nReviewer #3 (Remarks to the Author):\nI still have a lot of concerns about how the results are presented, particularly with the use of\nmicro-tensile testing. The values should not be compared with other sanples where macro-\ntesting is used. This gives a false impression that the properties are a lot better than they\nactually are. In SSupplementary Figure 2 we see the Mg14%NbC alloy has a yield strength\nof approximately 150MPa with an elongation to fracture of about 1%. This is typical for alloys\nsuch as this one. This is very different from the 678MPa yield strength improvement\nmentioned in the manuscript. If this is only compared with other micro tensile tests that\nwould be fine, but that does not seem to be the case.\nWe thank the reviewer for the careful evaluation of our revised manuscript and for raising\nthis important concern regarding the comparison of micro- and macro-tensile data.\nWe agree that elongation to fracture can depend on specimen size and that direct\ncomparison between micro- and macro-scale tensile tests may lead to misleading\ninterpretation. To address this concern, revised Fig. 2d has been added to the revised\nmanuscript. The updated figure now includes only literature data (see Supplementary Table\n3) obtained from comparable specimen sizes employing micro-tensile testing. The revised\nFig. 2d (Mg-NbC versus micro-tensile test data from the literature) shows the same overall\ntrend as that presented in the previous version of the figure.\nFigure 2b (inset) presents the yield strength increment (Δσ ) between the monolithic alloy\nYS\nand the corresponding reinforced alloy, where each comparison is made using the same\ntesting methodology within the respective study. Therefore, the comparison is\nmethodologically consistent and does not involve mixing micro- and macro-tensile data\nwithin a single Δσ calculation. Furthermore, yield strength is determined from the initial\nYS\nplastic deformation region and is not influenced by early fracture behaviour. The load cells\nused in both micro- and macro-tensile testing systems are calibrated, ensuring reliable force\nmeasurement and, consequently accurate stress determination. Therefore, the reported\nΔσ values reflect the intrinsic strengthening contribution of the reinforcement under the\nYS\nspecific testing conditions. For these reasons, Fig. 2b (inset) does not create a misleading\nimpression and has been retained without modification."}, {"round": 3, "reviewer_comments": {"reviewer_3": "Whilst this is an improvement, I request that the authors specifically say that the tensile results are from micro-tensile tests\nwithin the caption in Fig 2a. That is:\na. 'Micro'-tensile stress-strain curves\nThis will make it clear to the readers what they are looking at."}, "author_response": "Response to Reviewer #3’s Comments (NCOMMS-25-78483B)\nReviewer #3 (Remarks to the Author):\nWhilst this is an improvement, I request that the authors specifically say that the tensile\nresults are from micro-tensile tests within the caption in Fig 2a. That is:\na. 'Micro'-tensile stress-strain curves\nThis will make it clear to the readers what they are looking at.\nResponse: The word “micro” has been added to the caption of Fig. 2a, as suggested by the\nreviewer, to clarify that the tensile results are obtained from micro-tensile tests."}]

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