Achieving comprehensive performance regulation in Cu matrix composites
via CNT modified design
Yuansen Chen a,b, Jinyang Cao a,b, Wenmin Zhao a,b,c,d,*, Yikun Li a,b,
Baixiong Liu a,b,c,**, Hui Guo a,b,c, Qiuwen Liu a,b,c
a State Key Laboratory of Nonferrous Structural Materials, Jiangxi University of Science and Technology, Ganzhou, 341000, China
b School of Materials Science and Engineering, Jiangxi University of Science and Technology, Ganzhou, 341000, China
c Gannan Laboratory, Ganzhou, 341000, China
d Yunnan Key Laboratory of New Materials Preparation and Processing, Kunming, 650093, China
A R T I C L E I N F O
Keywords:
Cu matrix composites
Carbon nanotube
Carbonized polymer dot
Mechanical properties
Electrical conductivity
A B S T R A C T
To alleviate the trade-off between strength-conductivity in carbon nanotube (CNT)/Cu composites, by using a
one-step hydrothermal method to synthesize CPD in situ and simultaneously grafting it onto the surface of CNT, a
CPD@CNT composite reinforcement is successfully prepared. Subsequently, the corresponding composites are
fabricated by powder metallurgy. The results show that the CPD@CNT prepared at 190 ◦C exhibits a significant
strengthening effect on the Cu matrix composites. Specifically, it has excellent tensile strength (389 MPa), good
electrical conductivity (91.64 % IACS), and outstanding thermal stability (softening temperature>510 ◦C). This
is attributed not only to the abundant oxygen-containing functional groups on the CPD surface, through which
the dispersibility of CNT is enhanced and the interface wettability between the reinforcement and the matrix is
improved, but also to the differential thermal stability of these functional groups, partial degradation occurs at
the interface, resulting in the formation of a Cu2O transition phase that ultimately achieves robust interface
bonding. Furthermore, significant dislocation strengthening is achieved via dislocation pile-up and the Orowan
mechanism, owing to CPD@CNT dispersed at grain boundaries and within grains. Concurrently, an efficient
electron transport pathway is established by the robust interface bonding, whereby interface electron scattering
is effectively suppressed. This study provide a novel strategy for the design of high-strength, high-electrical
conductivity Cu matrix composites.
- Introduction Pure Cu exhibits excellent electrical and thermal conductivity, but it has relatively low mechanical strength and is prone to softening at high temperatures, which severely limits its application in today’s complex and diverse scenarios. Therefore, by adding high-performance rein forcement to prepare Cu matrix composites, while retaining the high electrical conductivity of the Cu matrix and enhancing its mechanical strength and high-temperature stability, has become a research hotspot in this field. Cu matrix composites are widely acknowledged as pivotal materials in critical sectors, including electronic packaging, power transmission, and aerospace, by virtue of this excellent electrical and thermal conductivity coupled with favorable manufacturability. They are considered among the cornerstone materials underpinning modern electronic information and energy technology systems [1–4]. Carbon nanotube (CNT), as typical representatives of one-dimensional nano carbon materials, are considered ideal reinforcements for Cu matrix composites owing to it ultra-high tensile strength, excellent elastic modulus, and outstanding electrical properties [5–7]. Theoretically, The properties of CNT and the Cu matrix are highly complementary, and a composite system that integrates high strength with high electrical conductivity is expected to be constructed [8]. However, during prac tical fabrication, CNT tend to severely agglomerate due to van der Waals forces. Additionally, the poor physical wettability between CNT and the Cu matrix results in insufficient interface bonding strength [9,10]. Furthermore, due to the significant differences in their thermal expan sion coefficients, interface stress concentration and structural defects are caused [11]. Not only are the load transmission and carrier transport
- Corresponding author. State Key Laboratory of Nonferrous Structural Materials, Jiangxi University of Science and Technology, Ganzhou, 341000, China.
** Corresponding author. State Key Laboratory of Nonferrous Structural Materials, Jiangxi University of Science and Technology, Ganzhou, 341000, China.
E-mail addresses: zhaowenmin@jxust.edu.cn (W. Zhao), liu_micro@126.com (B. Liu).
Contents lists available at ScienceDirect
Ceramics International
journal homepage: www.elsevier.com/locate/ceramint
https://doi.org/10.1016/j.ceramint.2026.04.440
Received 23 February 2026; Received in revised form 9 April 2026; Accepted 30 April 2026
Ceramics International 52 (2026) 29145–29155 Available online 30 April 2026 0272-8842/© 2026 Published by Elsevier Ltd.
efficiency severely impaired by these issues, but their enhancement
potential is also constrained. At the same time, intense interface electron
scattering is triggered, ultimately leading to a significant deterioration
in the performance of the composites [12–14]. Consequently, the prac
tical application of CNT/Cu composites is severely restricted by these
factors. Therefore, developing effective interface and structure control
strategies to balance the competing performance bottlenecks of
“strength - conductivity” has become an urgent research priority in this
field.
In recent years, many studies have focused on optimizing the struc
ture and designing the functions of CNT through multi-dimensional
synergistic strengthening strategies, and introducing secondary phases
to construct composite reinforcement systems. Furthermore, interfacial
synergistic effects and multi-mechanism coupling effects must be fully
exploited to surmount the performance limitations inherent to single
reinforcement phases, thus the comprehensive properties of composites
to be markedly enhanced. For example, Feng et al. [15] fabricated Cu
matrix composites reinforced with titanium carbide (TiC)-modified CNT
and in-situ nano-TiC particles. By taking advantage of the respective
merits of 1-D CNT and 0-D TiC, the simultaneous enhancement of
strength and ductility are achieved in the Cu matrix composites. Luo
et al. [16] coated the surface of CNT with Cu to promote the dispersion
of CNT and enhance the interface bonding between CNT and Cu. This
resulted in a more uniform distribution of strain in the composites,
thereby the strength improved from 14% to 26.5%. Chen et al. [17]
treated CNT with acid and uniformly deposited a Cu (or Ni) layer on the
surface of CNT via the electroless deposition method. Compression tests
demonstrated that the composite containing 2 vol% Ni-coated CNT
exhibited a maximum compressive yield strength of 554 MPa, which is
55% and 30% higher than that of pure Cu and the composite reinforced
with have significant advantages, the inherent properties of the CNT are
sacrificed, and the operation process is also considered to be rather
complex. Furthermore, not only are substantial economic costs incurred
by the acids and metal salts required for these experiments, but signif
icant environmental risks and challenges are also posed by waste solu
tion treatment. As a novel zero-dimensional carbon material with a
core-shell structure, carbonized polymer dot (CPD) are composed of
layers of graphene-like microstructures in the core, which are stacked
together. The shells of these CPD are modified with abundant
oxygen-containing functional groups. As a result, it possess the ability to
uniformly disperse in polar solvents and exhibit excellent electronic
transport properties [18,19]. However, when using CPD as the rein
forcement, the structure of CPD is inherently unstable and its single
reinforcing effect is limited. Based on these properties, the CPD
employed to improve the dispersion of CNT in the Cu matrix presents a
feasible strategy.
Consequently, in this work, a one-step hydrothermal method was
adopted to synthesize CPD, and it was simultaneously fixed on the
surface of CNT. This not only significantly shortened the time and eco
nomic cost required for acid treatment of CNT, but also effectively
improved the efficiency. Subsequently, high-strength and high-
conductivity CPD@CNT/Cu composites were fabricated via powder
metallurgy. The comprehensive performance of the CPD@CNT/Cu
composites was significantly enhanced by taking full advantage of the
synergistic performance of CPD and CNT. The hydrothermal reaction
temperature was precisely regulated to achieve the optimal conditions
for the anchoring and combination of CPD on CNT. This study provides
an important technical approach and research insight for the prepara
tion of Cu matrix composites with excellent dispersibility, strong inter
face bonding and outstanding comprehensive properties, thereby
offering a novel design concept and technical route for the development
of high-performance Cu matrix composites with integrated structural
and functional properties.
2. Experimental methods
The raw materials used in the experiment included: Hydroxyl-
functionalized multi-walled carbon nanotube (CNT, diameter: 10–20
nm, length: 0.5–2 μm, purity: 95%, Nanjing XFNANO Materials Tech.
Co., Ltd., China), Citric Acid (C6H8O7, CA, 99.5%, Shanghai Macklin
Biochemical Technology Co., Ltd., China), Ethylenediamine (C2H8N2,
EDA, 98%, Shanghai Macklin Biochemical Technology Co., Ltd., China),
Fig. 1. Schematic illustration of the preparation process of (a) CPD@CNT composite powder; (b) CPD@CNT/Cu composites.
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and electrolytic Cu powder (particle size: 325 mesh, purity: 99.9%,
Shanghai Naiou Nano Technology Co., Ltd., China).
2.1. Synthesis of CPD@CNT and composites
As illustrated in Fig. 1(a), a one-step hydrothermal method was
adopted to synthesize CPD and graft it onto CNT, yielding the CPD@CNT
hybrid. Specifically, CNT were first ultrasonically dispersed in deionized
water, followed by the sequential addition of CA and EDA. The mixture
was subjected to ultrasonic treatment to form a homogeneous suspen
sion, which was then transferred into a supercritical reaction autoclave
for hydrothermal reaction under continuous stirring. After dialysis and
freeze-drying, the CPD@CNT powder was obtained. To investigate the
effect of functionality degree on CNT, the synthesis reaction was con
ducted at four different temperatures (150 ◦C, 170 ◦C, 190 ◦C, 210 ◦C).
The resulting products were sequentially labeled as Samples A, B, C and
D to facilitate their identification and differentiation in subsequent
studies.
As shown in Fig. 1(b), the CPD@CNT composite powder with a mass
fraction of 0.3% was mixed with pure Cu powder in an ethanol medium
via variable-speed ball milling (200 rpm, 8 h + 400 rpm, 5 h). Subse
quently, the milled powder was reduced in a hydrogen atmosphere and
consolidated by spark plasma sintering (SPS). The sintered bulk was
subsequently subjected to multi-pass hot rolling with a total deformation
of approximately 65%, and the CPD@CNT/Cu composites was ulti
mately fabricated.
2.2. Materials characterization
The phase composition of the composite powders was determined
using X-ray diffraction (XRD, Shimadzu XRD-6100). The powder
morphology and composite microstructure were examined by scanning
electron microscopy (SEM, MLA650F). The morphology of the rein
forcement and the interfacial structure of the composite were charac
terized by transmission electron microscopy (TEM, Tecnai G2 F20).
Raman spectroscopy (Horiba LabRAM HR Evolution) was utilized to
investigate the carbon structure of the CPD@CNT hybrid material.
Electron backscatter diffraction (EBSD, attached to the MLA650F SEM)
was used to evaluate the average grain size of the composite. The Vickers
hardness of the CPD@CNT/Cu composites was measured at room tem
perature using an HV-1000A tester under a 100 g load with a 10 s dwell
time, and the electrical conductivity was determined at room tempera
ture using an eddy current conductivity meter (Sigma 2008A). Finally,
the mechanical properties of the composites were evaluated with a
universal testing machine (UTM 4204) at a tensile rate of 0.1 mm/min.
3. Results and discussion
3.1. Microstructural characterization
The microstructure of the pristine CNT and the as-prepared
CPD@CNT composite powders is displayed in Fig. 2. The macroscopic
state of pristine CNT and Samples A–D, after they are ultrasonically
dispersed in ethanol and subsequently allowed to stand for 24 h under
UV irradiation, is illustrated in Fig. 2(a). The highest fluorescence in
tensity under UV light is exhibited by Sample C, indicating that the
Fig. 2. (a) Photos of the prepared CPD@CNT in UV light (upper) and daylight (bottom), respectively; (b) TEM and HRTEM images (insets) of original CNT; (c)
sample A; (d) sample B; (e) sample C; (f) sample D.
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formation of CPD is rendered most effective under this temperature
condition. After standing, the pristine CNT are completely agglomerated
and settled, whereas the CPD@CNT composite powders remain uni
formly dispersed, which demonstrates excellent dispersion stability. A
TEM image of the pristine CNT is shown in Fig. 2(b), the corresponding
high-resolution TEM (HRTEM) image in the lower-right corner reveals
that pristine carbon nanotubes have a hollow tubular structure, with a
diameter of approximately 30 nm and a length of 0.5–2 μm, and their
walls are characterized by smoothness without obvious defects. How
ever, due to the strong van der Waals forces and high specific surface
area between CNT, they are prone to agglomeration [20]. TEM images of
the CPD@CNT composite powders for Samples A–D are presented in
Fig. 2(c–f), and the dispersion state of the CNT and the generation of
CPD on their surfaces are reflected by these images. With the increase in
hydrothermal reaction temperature, the agglomeration degree of CNT
first decreases and then intensifies, while the amount of CPD generated
on their surfaces also shows a trend of initial increase followed by
subsequent decrease. This indicates that a significant influence is exer
ted by the reaction temperature exerts on the surface structure of CNT
and the nucleation behavior of CPD. After CPD are successfully gener
ated on the surface of CNT, the dispersion stability of CNT in solvent is
improved by the abundant oxygen-containing functional groups on CPD.
The upper-right and lower-right insets in Fig. 2(e) are HRTEM images,
lattice fringes with interplanar spacings of 0.35 nm and 0.21 nm are
displayed by these images, which correspond to the (002) plane of the
CNT and the (100) plane of the CPD deposited on their surface,
respectively. This further confirms that CPD are successfully formed on
the surfaces of CNT.
Distinct diffraction peaks at 2θ = 26.14◦are indicated by XRD
diffraction analysis (Fig. 3(a)) for all CPD@CNT composite powders, and
these peaks are assigned to the (002) crystal plane of CNT. It is note
worthy that variations in the diffraction peak profiles of Sample A to D
are shown with the change of hydrothermal reaction temperature:
relatively broadened peaks are presented by Sample A and D, whereas
the sharpest and most intense diffraction characteristics are demon
strated by Sample C. This phenomenon indicates that an enhanced de
gree of graphitization and a more perfect crystal structure are achieved
for Sample C, from which it can be inferred that this specific processing
parameter is most favorable for the improvement of graphitization in
CPD@CNT composite reinforcement.
Fourier Transform Infrared (FTIR) spectroscopy is utilized to sys
tematically analyze the surface functional groups of the CNT. As shown
in Fig. 3(b), the presence of O-H (~3420 cm−1), C-H (~2930 cm−1),
C=O (~1700 cm−1), C=C (~1580 cm−1), N-H (~1380 cm−1), and C-N
Fig. 3. Microstructural characterization of CPD@CNT: (a)XRD; (b) Infrared spectrum; (c) Raman spectrum.
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(~1210 cm−1) is indicated by the FTIR spectra. It is noteworthy that the
C-N bond is clearly visible only in Sample A and B, whereas in Sample C
and D, it is significantly weakened or even disappears. The reason for
this phenomenon may be that at relatively low hydrothermal reaction
temperatures, the synthesis of CPD is primarily dominated by poly
merization and dehydration reactions between the precursors CA and
EDA. During this process, a “polymer framework” mainly composed of
polyamide-like structures is formed, accompanied by initial carboniza
tion. As a result, a large number of amide bonds (-CO-NH-) are retained
by the CPD, which are manifested as strong C-N and C=O bonds in the
FTIR spectra [21]. On the other hand, explains why the C=O peak in the
FTIR spectra of Sample A and B has a higher intensity, while it is
weakened in Sample C and D. When the hydrothermal reaction tem
perature is further increased, the degree of carbonization is intensified,
which leads to the pyrolysis of unstable C-N bonds. Some heteroatoms
are escaped in the form of volatile small molecules, while structural
rearrangement is undergone by the remaining carbon atoms, and larger
and more ordered sp2 structures are formed accordingly. Charge trans
port scattering is thus greatly reduced, and the electrical conductivity of
the material is further enhanced. It is confirmed by the FTIR results that
CPD are successfully grafted onto the surfaces of CNT. The polar func
tional groups contained in them not only enhance the stability and hy
drophilicity of the composite powder in polar solvents, so that uniform
dispersion of the powder in the matrix is promoted, but also facilitate the
improvement of interface wettability between the CNT and the Cu ma
trix, such that effective pathways for interface electron transport are
provided [22].
The defect density and structural ordering of carbon materials are
effectively reflected by the intensity ratio of ID/IG in Raman spectros
copy [23]. As shown in Fig. 3(c), the ID/IG ratios for Sample A to D are
0.297, 0.322, 0.386, and 0.616, respectively, and a clear monotonic
increasing trend is exhibited by these ratios. From XRD analysis, the
sharpness of the characteristic peaks is interpreted as a reflection of the
overall crystalline integrity of the material, with a typical detection
depth in the micrometer range. In contrast, the surface and local states of
the material are primarily probed by Raman spectroscopy, with a typical
analysis depth ranging from several to tens of nanometers. With the
increase in hydrothermal temperature, the amount of CPD generated on
the surfaces of CNT is continuously increased. Concurrently, a signifi
cant rise in the intensity of the D-band in Raman spectra is collectively
led to by abundant surface functional groups and new structural defects
formed between the CPD and CNT, which is manifested as a continuous
increase in the ID/IG ratio. The discrepancy between the XRD and Raman
results is precisely demonstrated to indicate that CPD are successfully
nucleate and grown on the CNT surfaces without significantly compro
mise to the overall skeletal integrity of the CNT. A sudden increase in the
values observed for Sample D indicates that excessively high tempera
tures also cause certain damage to the CNT structure itself.
The distribution state of the reinforcement phase in the CPD@CNT/
Cu composite powder is systematically characterized by SEM. As shown
in Fig. 4(a–d), the CPD@CNT reinforcements are uniformly distributed
throughout the composite powder as a whole. However, a small number
of clusters formed by the entanglement of CPD@CNT are still observed
in Sample A (see Fig. 4(a)). With the increase of hydrothermal temper
ature, the agglomeration phenomenon is significantly alleviated in
Sample B (see Fig. 4(b)). For Sample C (see Fig. 4(c)), the CPD@CNT
reinforcements are evenly distributed without any obvious agglomera
tion. This indicates that under these condition, the nucleation effect of
CPD on the surface of CNT is the most optimal. Thus, the surface func
tional groups of CPD can more effectively promote the dispersion of
CNT. Meanwhile, bridge-like and penetrating structures are observed,
which demonstrates that close physical contact and geometric
Fig. 4. Microstructure of CPD@CNT/Cu composite powders: (a) Sample A; (b) Sample B; (c) Sample C; (d) Sample D.
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compatibility are established between the reinforcements and the Cu
powder particles, rather than simple physical adsorption. Nevertheless,
when the temperature is further increased to the condition corre
sponding to Sample D, local agglomeration reoccurs in the composite
powder. It is inferred that excessively high hydrothermal temperature
causes structural damage to the CNT themselves, the nucleation sites of
CPD is reduced, and the surface modification and dispersion-promoting
effects of CPD are weakened. This observation is consistent with the
conclusions drawn from the TEM morphology and Raman spectrum
analysis of the aforementioned CPD@CNT composite powder, which
collectively verifies the significant influence of hydrothermal tempera
ture on the dispersion behavior of the reinforcement phase.
TEM and HRTEM are employed to analyze the interfacial micro
structure of Sample C in Fig. 5. It is shown in Fig. 5(a) that CPD@CNT
are not only present at grain boundaries but also distributed within the
Cu matrix. A clean interface without obvious gaps or pores is revealed by
magnified observation of a CPD@CNT at a grain boundary (Fig. 5(b)).
The interface bonding between CPD@CNT and the Cu matrix is further
subjected to high-resolution analysis (see Fig. 5(c)). In Fig. 5(d), a
distinct transition zone of Cu2O is observed between the CPD@CNT and
the Cu matrix. This zone is attributed to the decomposition of oxygen-
containing functional groups on the CPD during the SPS sintering pro
cess, by which active oxygen atoms are released. These active oxygen
atoms are then reacted with Cu atoms from the Cu matrix, leading to the
formation of a Cu2O transition zone at the interface. This observation is
consistent with related studies in which oxygen-containing functional
groups are utilized to improve the interfacial bonding of CNT/Cu com
posites [22,24,25]. The FFT and IFFT images of the region within the
white box are shown in Fig. 5(e1, e2). The observed interplanar spacings
of 0.209 nm and 0.213 nm are assigned to the Cu (111) plane and the
Cu2O (020) plane, respectively. A lattice mismatch of 10.57% is calcu
lated. The clear lattice fringes observed indicate the formation of a
semi-coherent interface. Lattice strain accommodation at the interface is
achieved by the generation of dislocations, which is beneficial for the
formation of a strong and tough interface bond. [26]. The Cu2O nano
particles act as a bridge between CNT and the Cu matrix, and the
Fig. 5. TEM images of CPD@CNT/Cu composite of sample C; (a-b) TEM images of CPD@CNT/Cu; (c) enlarged TEM image of the white box in b; (d) enlarged TEM
image of the white box in c; (e1-e2) FFT and IFFT of the white box in d; (f) enlarged TEM image of the white box in c; (g1-g2) FFT and IFFT of the white box in f.
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effective enhancement of interfacial wettability as well as the reducing
of interfacial energy are both considered to be associated with this. In
Fig. 5(f), direct and tight bonding between the CPD@CNT and the Cu2O
interface is observed. Based on the corresponding FFT and IFFT images,
lattice spacings of ~0.206 nm and ~0.2135 nm are identified, corre
sponding to the CPD (1 12) and Cu2O (00 2) crystal planes, respectively.
This interfacial relationship collectively contributes to a continuous,
Fig. 6. Formation of the transition zone.
Fig. 7. (a-c) Different dislocation types of sample C; (d) IFFT image; (e) Inverse pole Fig (IPF) of the EBSD image; (f) Grain size statistics chart.
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gradient-strengthened interface structure extending from the CNT,
through the Cu2O transition layer, to the Cu matrix.
As shown in Fig. 6, the formation process of Cu2O is directly visu
alized. During the sintering process, oxygen atoms are released from the
oxygen-containing functional groups on the CPD surface, and a tighter
transition layer is constructed between the CNT and Cu. The weak
physical contact relying on van der Waals forces between CNT and Cu is
upgraded to strong chemical bonding achieved through C-O-Cu link
ages. In terms of mechanical properties, this transition layer acts as
uniformly distributed nano-rivets, enabling efficient stress transfer from
the Cu matrix to the CNT, which provides a fundamental guarantee for
the remarkable strength enhancement of the material. In terms of
electrical properties, due to the low content and dense structure of Cu2O,
it mainly serves as a bridging connection, and its influence on the overall
electron transport is minimized.
Typical dislocation configurations observed in the CPD@CNT/Cu
composites are displayed in Fig. 7. As is shown in Fig. 7(a), significant
dislocation pile-ups are observed near grain boundaries. It is indicated
that the transmission of dislocations into adjacent grains is effectively
impeded by grain boundaries, leading to the generation of localized
stress concentration. When mobile dislocations encounter CPD@CNT
reinforcements, two distinct interaction mechanisms are observed. The
first is illustrated in Fig. 7(b), where a large number of dislocations
gliding on the same slip system are strongly pinned upon encountering a
Fig. 8. (a) Bar chart of relative conductivity and hardness; (b) Engineering stress-strain curve; (c) True stress-strain curve; (d) Hardening rate curve; (e) Comparison
of tensile and electrical conductivity properties with previously reported Cu composites [5,9,28–36].
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CPD@CNT, and a zone of high-density dislocation accumulation is
formed around it. This phenomenon is primarily attributed to the
mismatch in the coefficients of thermal expansion between the CNT and
the Cu matrix [27]. In the second mechanism, when a mobile dislocation
meets an impenetrable CPD@CNT, the dislocation line is not simply
piled up but can be bowed around and bypassed by the dislocation line,
as shown in Fig. 7(c). It is revealed from the cross-sectional view that the
dislocation line is circumvented by the reinforcement and its continuous
motion is maintained, while a residual dislocation loop encircling the
reinforcement is simultaneously left behind. These dislocation loops and
the surrounding high-density dislocation network are clearly resolved at
the atomic scale by HRTEM (Fig. 7(d)). The key microscopic mechanism
enabling the composites to achieve high strength and high
work-hardening capability is constituted by these multi-scale disloca
tion interactions, which range from grain boundaries to grain interiors
and from pinning to bypassing mechanisms. It is indicated by the IPF
analysis of the EBSD image in Fig. 7(e) that a typical equiaxed grain
structure with random grain orientation distribution and no significant
texture formation is exhibited by the CPD@CNT/Cu composite. The
average grain size of approximately 1.83 μm is obtained from the cor
responding grain size statistical distribution (see Fig. 7(f)). This grain
size is found to be significantly smaller than that of conventional pure Cu
or single-reinforcement composites, and the notable grain refinement
effect of the CPD@CNT composite reinforcement on the Cu matrix is
directly confirmed. This refinement is primarily attributed to the
pinning effect of the reinforcements on grain boundary migration during
sintering and hot rolling, by which grain growth is effectively inhibited.
The refined equiaxed grain structure is shown to contribute significantly
to strengthening via the Hall-Petch mechanism, while the coordination
of multiple slip systems is also facilitated, thereby enhancing the uni
form deformation capability of the material. This microstructural
characteristic,
combined
with
the
aforementioned
dislocation
strengthening mechanisms, collectively forms the microstructural
foundation for the composite to achieve high strength and good
ductility.
3.2. Electrical conductivity and mechanical properties
Bar charts of the electrical conductivity and hardness values of the
composites are presented in Fig. 8(a). The overall variations are rela
tively minor, and high electrical conductivity (90.18%–91.64% IACS) is
maintained by all samples. The preservation of conductivity can be
attributed to the intrinsic sp2-carbon character and conductivity of the
CPD themselves, coupled with the positive role played by them in
improving the CNT-Cu interface and reducing electron scattering.
The engineering stress-strain curves of the CPD@CNT/Cu compos
ites are shown in Fig. 8(b). A marked enhancement in tensile properties
is observed for Sample C, with an ultimate tensile strength (UTS) of
389.82 MPa and a yield strength (YS) of 331.23 MPa. This is attributed
to the formation of a Cu2O transition zone between CPD@CNT and the
Cu matrix. Efficient load transfer is enabled by this type of interface via
strong chemical bonding, while electron scattering is simultaneously
minimized due to its clean and well-defined nature. The enhancement in
strength demonstrates that the introduction of CPD enables the
achievement of significant strengthening while successfully circum
venting the dilemma of a sharp decline in electrical conductivity
induced by traditional reinforcement phases. However, a significant
enhancement in strength is accompanied by a moderate decrease in
plasticity. The fracture elongation of Sample C (13.32%) is found to be
lower than that of Sample B (21.73%), where optimal ductility is
exhibited. This observation is attributed to the fact that excessively high
strength renders dislocation motion extremely difficult, and the critical
stress for microcrack nucleation is reduced accordingly [15].
The true stress-strain and work hardening rate curves of the com
posites are shown in Fig. 8(c) and (d), respectively. The work hardening
rate curves clearly reveal the intrinsic differences in the deformation
behavior between the CPD@CNT/Cu composites. A very high initial
work hardening rate is observed for Sample C, which, however, is
decreased drastically with increasing strain. It is indicated that the
reinforcement (CPD@CNT) within this sample was extremely effective
at hindering dislocation motion in the early stages of deformation.
However, swift exhaustion of its strain hardening capacity is induced by
the rapid accumulation of dislocations, thereby limiting its plastic
deformability. In sharp contrast, although a slightly lower initial work
hardening rate is exhibited by Sample B, its curve displays the most
gradual decline. Excellent and stable work hardening capability is
maintained throughout the entire deformation process. This sustained
strengthening effect enables effective resistance to localized necking,
thus achieving the maximum uniform plastic deformation.
Good electrical conductivity is exhibited by Sample C while consid
erable tensile strength is possessed. Although its elongation at break is
decreased compared with that of Sample B, which exhibits optimal
plasticity, this moderate sacrifice of ductility is considered acceptable in
the design of multifunctional structural materials, and its absolute value
is found to be significantly higher than that of many reported advanced
Cu matrix composites [5,9,28–36]., as comparatively illustrated in Fig. 8
(e).
The hardness evolution curve of Sample C after annealing for 1 h at
temperatures ranging from 300 to 800 ◦C is displayed in Fig. 9. It is
observed that as the annealing temperature increases, the hardness is
first decreased slowly, then dropped sharply, after which the rate of
decrease is moderated again. This corresponds to distinct microstruc
tural evolution stages: the recovery stage, in which internal stresses are
relieved without alteration of the deformed microstructure. The
recrystallization stage, in which new, strain-free grains are nucleated
and grown. And the grain growth stage, in which grains are grown to
reduce the overall system energy [37]. According to the Chinese Na
tional Standard (GB/T 33370–2016), the softening temperature of a
specimen is defined as the annealing temperature at which its micro
hardness is decreased to 80% of the original value. The material’s sta
bility and tolerance under high-temperature conditions are reflected by
the softening temperature, and it is served as a key indicator for eval
uating the high-temperature resistance of Cu. As is indicated in Fig. 8,
the softening temperature of sample C is greater than 510 ◦C, with a
corresponding hardness of 90.448 HV at that point. This softening
temperature is found to be significantly higher than the typical recrys
tallization temperature range of pure Cu (200–300 ◦C), indicating that
the high-temperature softening resistance of the composite is substan
tially enhanced by the introduction of CPD@CNT. This improvement is
primarily attributed to the synergistic pinning effect of the
Fig. 9. High-temperature softening resistance curve.
Y. Chen et al.
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multi-dimensional reinforcements. Grain boundary migration is effec
tively hindered by CPD@CNT distributed at grain boundaries, by which
nucleation and growth during recrystallization are suppressed. Mean
while, dislocations are subjected to stable obstacles by those dispersed
within the grains, delaying dislocation recovery and rearrangement at
elevated temperatures. The temperature at which significant softening
occurs is collectively elevated by approximately 210 ◦C due to this
multi-scale pinning effect at both interfaces and defects, fully demon
strating that the core advantage of nano-carbon reinforcements is real
ized in enhancing the thermal stability of Cu matrix composites.
4. High-strength, high-conductivity mechanism
To achieve high strength and high electrical conductivity in Cu
matrix composites, the aim is to simultaneously meet the load-bearing
requirements of these materials as structural components and the effi
cient energy transmission demands of them as electrical conductors.
The multi-scale strengthening and conductive scattering transport
pathways constructed by CPD@CNT within the Cu matrix are illustrated
in Fig. 10. Its strengthening mechanism is manifested as dual dislocation
obstruction at grain boundaries and within grains. As is shown in Fig. 10
(a), rigid non-deformable obstacles are acted as by CPD@CNT located at
grain boundaries, leading to the accumulation of a large number of
dislocations at their frontiers. The slip of dislocations across grain
boundaries to adjacent grains is thereby effectively hindered, which
significantly enhances the grain boundary strengthening effect of the
composites [38]. According to the dislocation pile-up theory, a strong
back stress is generated by the accumulation of dislocations in front of
grain boundaries, thereby remarkably increasing the yield strength of
the material. The contribution of grain boundary strengthening is thus
constituted. Within grains, when a dislocation line moves on the slip
plane and encounters CPD@CNT, the dislocation line is bent and grad
ually bypasses the reinforcement phase, as direct cutting through the
latter is not permitted. Subsequently, driven by the sustained applied
stress, the bent dislocation line continues to bypass the reinforcement
phase. Owing to the considerable length of CPD@CNT, the dislocation
line proceeds to move forward after bypassing the reinforcement phase,
a behavior analogous to the Orowan bypass mechanism. This process not
only directly elevates the yield strength but also enables the continuous
multiplication of residual dislocation loops, which in turn endows the
material with excellent work-hardening capacity from another aspect.
At the same time, CPD@CNT, through its unique “candied haw
thorn” structure, not only greatly reduces electron scattering but also
constructs efficient charge transport channels, thereby ensuring high
conductivity. The core mechanism is found to lie in the fact that excel
lent electrical conductivity is possessed by the sp2 carbon core of CPD
itself, with which an efficient hybrid charge transport channel is formed
together with CNT. A pathway with a lower energy barrier for electron
transport from Cu to CNT and then back to Cu is provided by the formed
transition layer, by which electron scattering at the interface is greatly
mitigated and interfacial contact resistance is thereby significantly
reduced (see Fig. 10(b)).
5. Conclusion
In this work, CPD@CNT/Cu composites were prepared through a
one-step hydrothermal method and SPS process. CNT are evenly
dispersed in the Cu matrix with the help of CPD. CPD@CNT achieves
dislocation strengthening through the accumulation of dislocations at
grain boundaries and within grains, as well as the Orowan mechanism.
Moreover, the abundant oxygen-containing functional groups on the
CPD surface improve interfacial wettability. At the same time, the
unique structure of CPD@CNT constructs efficient charge transport
channels, significantly reducing electron scattering, thereby enabling
the composite material to possess both excellent mechanical properties
and electrical conductivity.
CRediT authorship contribution statement
Yuansen Chen: Writing – original draft, Validation, Investigation.
Jinyang Cao: Methodology, Formal analysis, Data curation. Wenmin
Zhao: Writing – review & editing, Supervision, Resources, Project
administration, Investigation. Yikun Li: Methodology, Investigation,
Formal analysis. Baixiong Liu: Project administration, Methodology,
Investigation. Hui Guo: Supervision, Investigation. Qiuwen Liu: Vali
dation, Investigation.
Declaration of competing interest
The authors declare that they have no known competing financial
interests or personal relationships that could have appeared to influence
the work reported in this paper.
Acknowledgment
This work was supported by the National Science Foundation of
China (Grant No. 52304382), Yunnan Key Laboratory of New Materials
Preparation and Processing (2024KF08), the Program of Jiangxi Pro
vincial Laboratory of Advanced Metal Materials and Critical Minerals
Development (2025GNJC02), Advanced Materials-National Science and
Technology Major project (2025ZD0619700), and the Pioneer Research
Program of Jiangxi Provincial Center for Interdisciplinary Studies in
Future-oriented Materials.
Data availability
Data will be made available on request.
Fig. 10. Strengthening mechanism diagrams: (a) Dislocation strengthening mechanism; (d) Conductive strengthening mechanism.
Y. Chen et al.
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