paper19-advmat-2026-bct-cu-nanowires-co2rr-c2plus

RESEARCH ARTICLE

Stabilization of Unconventional Body-Centered Tetragonal Phase in Copper Nanowires for Efficient Carbon Dioxide Electroreduction to Multi-Carbon Products Guozhi Wang1,2 Yangbo Ma1 Mingzi Sun1 Fengkun Hao1 Qianhui Wei3,4 Yunhao Wang1 1 1,2 1,2 1,2 1,2 Fu Liu Xiang Meng Liang Guo Mingzheng Shao Chaohui Wang Shuheng Hao1 Yuecheng Xiong1 Yanwei Lum5 Shengqi Chu6 Bolong Huang1 Zhanxi Fan1,2,7,8

Juan Wang1 Pengyi Lu1

1 Department of Chemistry, City University of Hong Kong, Kowloon, Hong Kong SAR, China 2 Hong Kong Branch of National Precious Metals Material Engineering Research Center (NPMM), City University of Hong Kong, Kowloon, Hong Kong SAR, China 3 GRINM (Guangdong) Institute for Advanced Materials and Technology, Foshan, China 4 State Key Laboratory of Advanced Materials for Intelligent Sensing, GRINM Group Co., Ltd., Beijing, China 5 Department of Chemical and Biomolecular Engineering, National University of Singapore, Singapore, Singapore 6 Beijing Synchrotron Radiation Facility, Institute of High Energy Physics, Chinese Academy of Sciences, Beijing, China 7 Hong Kong Institute for Clean Energy, City University of Hong Kong, Kowloon, Hong Kong SAR, China 8 City University of Hong Kong Shenzhen Research Institute, Shenzhen, China

Correspondence: Qianhui Wei (weiqianhui@grinm.com) Received: 15 March 2026

Revised: 18 April 2026

Bolong Huang (b.h@cityu.edu.hk)

Zhanxi Fan (zhanxi.fan@cityu.edu.hk)

Accepted: 18 May 2026

Keywords: carbon dioxide reduction reaction | crystal phase | electrocatalysis | metal nanomaterials | multi-carbon products

ABSTRACT Copper nanomaterials with the common face-centered cubic (fcc) phase have been widely used in the electrocatalytic carbon dioxide (CO2 ) reduction reaction (CO2 RR). However, copper with an unconventional phase is rarely reported as it is thermodynamically unfavorable. Here, through analyzing the strain within copper nanowires, we reveal the phase transition of copper from fcc to body-centered tetragonal (bct)/fcc heterophase. By systematically investigating copper nanowires with different diameters and copper nanocubes in CO2 RR, we explain the relationship between their crystal phase and catalytic performance. Compared with the standard fcc lattice, copper nanowires’ surfaces have different electron states due to a phase transition. Copper nanowires with a diameter of about 30 nm exhibit the optimum catalytic performance, and their Faradaic efficiency of multicarbon products is much higher than that of fcc copper nanocubes. Theoretical calculations have demonstrated that the presence of the strained bct phase induces significant upshifts of the d-band center, which not only improves the overall electroactivity but also optimize the C-C couplings, leading to improved Faradaic efficiency of multi-carbon products during CO2 RR.

Introduction

Electrocatalytic conversion of carbon dioxide (CO2 ) into valueadded chemicals or fuels, which can use clean energy such as photovoltaics, wind power, and hydropower as a driving force, holds great potential to curb carbon emissions and promote carbon neutrality [1]. In the electrochemical CO2 reduction reac-

tion (CO2 RR), CO2 can be converted to a series of single-carbon (C1 ) products (e.g., carbon monoxide (CO), methane (CH4 ), and formic acid (HCOOH)) and multi-carbon (C2+ ) products (e.g., ethylene (C2 H4 ), ethanol (C2 H5 OH), and n-propanol (C3 H7 OH)) [1–6]. Among these products, C2+ products have attracted more attention due to their higher economic value and larger energy density compared to C1 products [2, 7]. Copper (Cu)-based

Guozhi Wang, Yangbo Ma, and Mingzi Sun contributed equally to this work.

Advanced Materials, 2026; 38:e73600 https://doi.org/10.1002/adma.73600

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To date, great efforts have been devoted to finding ways to enhance the CO2 electroreduction to C2+ products on Cu-based catalysts [12–16]. Common types of catalysts include pure metallic Cu (nano)materials, Cu single-atom catalysts, Cu alloys, Cu heterostructures, Cu oxides, Cu-based metal organic frameworks, etc [10, 13, 17–23]. As for the pure metallic Cu (nano)materials, researchers mainly studied the effects of size, facet, defect, and morphology on catalytic performance [24–27]. Unfortunately, the crystal phase effect of Cu (nano)materials on CO2 RR has been almost neglected. This is because Cu tends to adopt the thermodynamically stable face-centered cubic (fcc) phase, and it remains challenging to modulate the crystal phase of Cu (nano)materials. Very recently, we have obtained unconventional hexagonal close-packed (hcp) Cu nanostructures through the epitaxial growth of Cu on gold templates, which show improved catalytic performance in CO2 RR compared with fcc Cu, indicating the paramount significance of crystal phase control [17, 28]. However, the unconventional phase of metallic Cu nanomaterials is almost limited to the hcp structure so far. In this work, we synthesize Cu nanowires (NWs) with different diameters (30 nm, 80 nm, and over 200 nm) and standard fcc phase Cu nanocubes (NCs), and the focused ion beam (FIB) technique is used to obtain the cross-section sample of individual Cu NWs. Then, the high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) images of the cross-section sample verified the distribution of defects and strains and phase transition from fcc to bct/fcc heterophase in Cu NWs. X-ray and ultraviolet photoelectron spectral analysis reveal their crystal structure and the electronic states of the surface and bulk phases. In electrocatalytic tests, the thinnest nanowires (bct/fcc Cu NWs-30) show the highest Faradaic efficiency (FE) about 80% for C2+ products at -1.15 V (vs reversible hydrogen electrode (RHE)). As the diameter increases, the FE of C2+ products gradually decreases, but the thickest nanowires’ (bct/fcc Cu NWs-200) performance is still better than that of fcc Cu NCs. In situ differential electrochemical mass spectrometry (DEMS) characterizations indicate that bct/fcc Cu NWs-30 decrease the onset potential for C2+ generation. Density functional theory (DFT) calculations suggest that the unconventional bct phase with strain effect is beneficial for the surface electroactivity of the Cu (100) surface, which not only promotes the adsorption of key reactants but also alleviates the barriers for C-C coupling and C2 H4 formation.

Results and Discussion

The understanding of the fivefold twinned structure has been evolving for more than half a century. The core of the problem is how the fivefold twin structure, which is composed of five regular tetrahedral subunits, fills the 7.35◦ solid angle deficiency. It began with Bagley’s homogeneous strain theory in 1965 [29]. Then R. de Wit proposed the disclination model in 1972 [30]. And later in 2008, Johnson et al. proposed a more reasonable explanation that 2 of 13

disclination and shear stress work together to fill the solid angle deficiency [31]. The unique fivefold twinned Cu decahedron and nanowire are shown in Figure 1a (Panels I-V). The decahedron is composed of five tetrahedron subunits, and if all five subunits maintain a regular tetrahedral structure, a solid angle deficiency will be generated (Panels I, II, and IV). In wet-chemical synthesis, the solid angle deficiency is spontaneously filled as the crystals grow (Panels II and III), which is the origin of their internal stress and strain. Analyzing stress and strain is an important way to understand phase transitions. Cu NWs could be regarded as the Cu decahedron growing along the fivefold axis ([110] crystal orientation), which also needs to fill the angle deficiency (Panels IV and V). Figure 1b shows a schematic cross-sectional view of five-fold twinned nanowires. The crystal orientation shown in Figure 1b applies to the bottom subunit.

2.1

Synthesis and Structural Characterization

By using different capping agents, we synthesized Cu nanowires with different diameters (Please see details in Supporting Information). During the synthesis, capping agents can selectively adsorb onto the Cu (100) facets and restrict their growth, thereby directing the axial growth of nanowires [32]. Different capping agents (e.g., hexadecylamine, octadecylamine, and oleylamine) will form molecular layers with varying packing densities upon adsorption on the (100) facets [33]. Consequently, Cu NWs with different diameters have been prepared by controlling the feeding ratios of metal precursors, reducing agents, and capping agents, as well as the reaction time and temperature (Figure S1). Here, the transmission electron microscopy (TEM) images show the Cu NWs with diameters of 30 nm, 80 nm, and over 200 nm. (Figure 1c; Figure S1). To study the internal stress and strain, we selected nanowires with diameters greater than 80 nm (Figure 1c) to prepare the cross-section sample using the FIB technique. To ensure successful cutting, we did not use nanowires with the smallest diameter. Figure 1d and Figure S2 show the TEM and scanning electron microscopy (SEM) images of the crosssection of the Cu NW. Before FIB cutting, Cu NWs were stored in ethanol at a low temperature. After cutting and before HAADFSTEM characterization, the sample was kept under vacuum to prevent oxidation. Five subunits and twin boundaries can be clearly seen in Figure 1d. The selected-area electron diffraction (SAED) pattern of the cross section (Figure S3) also shows the fivefold symmetry. Then, HAADF-STEM was used to observe the crystal structure in detail. In Figure 1e, the center of the Cu NW is brighter than the other area. Since the same element should have the same contrast, these light and dark variations are caused by high intensity of strain in the center. To facilitate our discussion of stress distribution later, here we use the crystal plane indices belonging to the FCC lattice for discussion. And we define that in each subunit, the direction from the center and perpendicular to the bottom edge is the vertical direction (red arrow in Figure 1e), and the direction perpendicular to this red arrow is the lateral direction (blue arrow in Figure 1e). Figure 1f shows a unit cell of the fcc phase, where the (110) facet represented by the red rectangle is the crystal plane where the cross section is located. Figure 1g shows the detail of the marked area by the blue rectangle in Figure 1e. It’s clear to observe that the atoms on both sides of the twin boundary (green line) are arranged symmetrically. The atoms in subunit 1 and subunit 5 Advanced Materials, 2026

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nanomaterials are a promising kind of catalysts that are able to facilitate the electroreduction of CO2 to C2+ products at ambient conditions [8]. However, Cu-based catalysts still suffer from unsatisfactory activity and selectivity toward C2+ products in CO2 electroreduction [1, 2, 9–11], especially for specific C2+ products.

both show a characteristic stacking sequence of “ABC,” and the twin boundary shows a stacking sequence of “ABA”. However, we found that the twin boundary is not always continuous from the center to the surface. Here exists a twin boundary migration phenomenon, which does not occur independently but many times. On the boundary between subunit 4 and subunit 5, there are 4 times of such twin boundary migration phenomena from the center to the surface (Figure 1h). Position 4 is used as an example to analyze this phenomenon (Figure 1i). The twin boundary here is not a coherent twin boundary (CTB). Because if we use either the green line as the boundary, the two sides are not completely symmetrical. As the twin boundary extends from the center to the surface, green line 1 stops at the white arrow, and green line 2 splits into two columns of atoms. After that, the new twin boundary moves down half an atomic plane. This grain boundary migration is not the bending of the grain boundary in Advanced Materials, 2026

one direction, but more like a Frank partial dislocation or an edge dislocation [34–40]. Normally, the migration is driven by shear stress, which verifies Johnson’s work [31, 37]. Twin boundary is a strong barrier to slip transmission, but in the disclination model, the circumferential stress (σθθ ) increases from the center to the surface, making dislocation across the CTB possible [37, 41, 42]. Thus, an additional atom in (110) plane could be regarded as a release of circumferential stress. Due to the migration of twin boundaries, the atomic arrangement at the white arrows is disrupted. In a dark-field STEM image, the contrast and clarity of the atoms around the migration area are reduced. Twin boundary migration affects 3–4 atomic layers along and perpendicular to the twin boundary direction ((111) facet), which could be a pinning phenomenon [38]. The details of positions 1–3 are shown in Figure S4a–c. The migration of grain boundaries is not always in the same direction. Positions 1 to 3 move up by 1.5, 1, and 2.5 3 of 13

FIGURE 1 Structural characterization of five-fold twinned Cu NWs. (a) Schematic drawing of a Cu decahedron and Cu NW with five-fold symmetry. The decahedron is composed of five tetrahedral subunits, and nanowires can be formed through the decahedron growth along the [110] direction. (b) Schematic drawing of the cross-section of Cu NWs and their state before and after suffering internal strain. The direction perpendicular to the cross section is the [110] zone axis. (c) TEM image of Cu NWs-80. (d) TEM image of a cross-section of Cu NWs. Five subunits are numbered as 1, 2, 3, 4, and 5. (e) HAADF-STEM image of the central part of the cross section in (d). The blue and red arrows define the lateral and vertical directions in subunit 1. (f) Schematic illustration of an fcc lattice. The red rectangle indicates the (110) facet, corresponding to the facet of the cross section. The black arrow indicates the [110] direction that is perpendicular to the cross section. (g) HAADF-STEM image of the selected area by the white rectangle in (e). (h) HAADF-STEM image of the twin boundary between subunits 4 and 5. (i) HAADF-STEM image of area 4 in (h). The green lines mark the twin boundary.

atomic planes, respectively. The schematic drawing of this twin boundary migration is shown in Figure S4d. The twin boundary migration produces multiple new atomic planes around the twin boundary. The multiple migrations release the lateral stress, and against it increases to higher levels [37, 38]. The migration of grain boundaries is random; not every grain boundary migrates (Figure S5, the twin boundary of subunits 2 and 3 doesn’t migrate). Here we propose that, under relatively high stress and large elastic strain, at certain locations, the twin boundaries are forced to break and migrate, forming dislocations in their original locations. New atomic planes are continuously added to (110) facets. This method effectively reduces the circumferential stress of the entire crystal and controls the strain within a smaller range of areas near the twin boundary. As circumferential stress exists in (110) facets and affects both lattice constants a and b, circumferential stress is released, that is, the stress in the a and b directions is released at the same time. To analyze the vertical and lateral stress distributions more intuitively, we directly measured the interplanar spacing at different positions. Owing to the limited measurement precision of the instrument, averaging over multiple atomic layers can effectively reduce the error and improve the measurement accuracy. Here, subunit 1 is taken as an example, and three areas (center, middle, 4 of 13

and bottom) are measured (Figure 2a,d,g). In the center part, considering that the high stress area is only in a small area around the center, not too many atomic planes are measured in vertical and lateral directions. In the 9 vertical atomic planes, the average interplanar distance is 1.80 Å (Figure 2b), which represents the interplanar spacing of (002) planes, and the constant of c is calculated to be 3.60 Å. The average interplanar spacing of the 19 lateral atomic planes is 1.30 Å (Figure 2c), which represents the interplanar spacing of the (220) facet. Since the cross section is not the (100) or (010) plane of the fcc lattice, lattice constants a or b cannot be directly measured. Based on this value, the average value of a and b can be calculated as 3.67 Å (All related values are summarized in Table S1). The spacing in the lateral direction is larger than in the vertical direction. The fast Fourier transform (FFT) pattern of the center part is shown in Figure S6. Based on −

−

the d-spacings of (002), (002), (220), and (220), the interplanar spacings of (001) and (110) can also be calculated as 3.61 Å and 2.60 Å, while those of (100) and (010) are 3.67 Å. The calculated values obtained from FFT can be regarded as the average value of the interplanar spacing in this region. Similarly, a and b are larger than c. When the regions used for FFT are too small, the pixels of the FFT result are limited, which could introduce errors. Therefore, we mainly use direct measurement results and use FFT results as auxiliary verification. The detailed information of Advanced Materials, 2026

FIGURE 2 Interplanar spacing analysis of subunit 1. (a) HAADF-STEM image of the center part. The red and blue rectangles represent the vertical and lateral directions, respectively. Vertical (b) and lateral (c) interplanar spacings of the center part. (d) HAADF-STEM image of the middle part. Vertical (e) and lateral (f) interplanar spacings of the middle part. (g) HAADF-STEM image of the bottom part. Vertical (h) and lateral (i) interplanar spacings of the bottom part.

X-ray and UV Photoelectron Spectral 2.2 Analysis There has been controversy over whether five-fold twinned nanowires adopt the bct phase or the body-centered orthorhombic (bco) phase [41, 43]. Here, X-ray diffraction (XRD) was used to investigate the structure of Cu NWs with different diameters (Figure 3a–e; Figure S9). Considering that copper light sources (λ = 1.54056 Å) usually have Kα and Kβ , Kβ filter was utilized to ensure a monochromatic light source [44]. Kα contains Kα1 and Kα2 , and their values were recorded by the equipment as 1.54059 and 1.54441 Å, respectively. Usually, the intensity ratio of these two sources is close to 2/1. Then Kα2 is removed during data processing. The peak intensity is normalized based on the intensity of (111) peak. The bct/fcc Cu NWs-30 show obvious side Advanced Materials, 2026

peaks at (200) and (220) peaks (Figure 3c,d). In Figure 3d, the (220) peak has one sharp peak and two weak side peaks. The sharp peak represents the unstressed Cu lattices. Its intensity is much stronger than the other two peaks. These unstressed lattices widely exist in the middle area of each subunit. The lowangle side peak (73.51◦ ) represents (220) facet in the center area. Because, based on the analysis in Figure 2, the lateral direction in the center and bottom parts shows tensile strain. The higher angle side peak (74.26◦ ) represents the stressed (202) and (022) facets at the center and bottom areas of each subunit. In this area, the vertical direction exhibits compressive strain. In Figure S9a, bct/fcc Cu NWs-80 have similar XRD results to bct/fcc Cu NWs-30. This result is not completely consistent with the bco model. In the reported bco models, the tensile strain increases gradually from the center to the surface [42]. And a ≠ b ≠ c, a and b suffer from different levels of tensile strain and will have two different intensity peaks at a lower angle, while c suffers from compressive strain and will create a peak at a higher angle [41, 45]. This means that all peaks are shifted and there is no sharp main peak. Here, the main peak in Figure 3d of bct/fcc Cu NWs-30 is in the same position as the (220) peak of fcc Cu NCs (Figures S9c and S10). This means that there are two kinds of bct phases in the nanowires. One is near the center, and the other is near the surface, a = b, and they are equally stretched, while c is compressed. That explains those 3 peaks in Figure 3d. In Figure 3c, the side peak is (002), and its maximum extends to 51.50◦ . Based on this angle and Bragg’s law λ = 2dsinθ, c under the highest compression is 3.55 Å, while the lattice constant corresponding to the main peak is 3.62 Å, suggesting the compression is 1.93% (The absolute value of the lattice constant measured by XRD and HAADF-STEM is different, which results from the systematic error of the equipment). This value in Figure S9a is 1.34%, which is higher than the 0.82% tested in Figure 2a, and also indicates that more pronounced compression may exist in nanowires with smaller diameters. Figure 3c shows one obvious side peak, and the left one is not obvious. This may be because its peak intensity is too weak and it is covered by the main peak or gets lost in the fluctuations of the baseline. As the diameter of the nanowire increases, its XRD pattern is very close to pure fcc Cu (Figure S9b,c). This may result from the ratio of the stressed area being too small compared with the stress-free areas in the nanowires, and the main peak is too high and thus covers these stressed peaks. The (311) peak of bct/fcc Cu NWs-30 (Figure 3e) has no obvious side peak. But in bct/fcc Cu NWs-80 (Figure S9a), when the main peak’s intensity is also very low, the side peak appears. In short, the phase transition in Cu NWs is from fcc to bct/fcc heterophase (Figure 3f,g), and with the increase of diameter, the Cu NWs’ compression and stretching phenomenon in XRD becomes less obvious. And reproducibility tests of samples synthesized in different batches further confirm that the side peaks in the XRD patterns do not arise from random errors or instrumental systematic errors, but are an intrinsic characteristic of Cu nanowires themselves (Figure S11). Besides, no peaks of CuO or Cu2 O were observed, proving that Cu NWs are in a metallic state, which is consistent with the X-ray photoelectron spectroscopy (XPS) results (Figure S12). The internal stress of Cu NWs may modify their surface electronic state and thus affect the catalytic performance. To verify this point, ultraviolet photoelectron spectroscopy (UPS) was used to further check the electronic state of different Cu NWs 5 of 13

the middle area is shown in Figure 2d–f. In Figure 2d, results can be obtained for any vertical and lateral atomic planes. No matter whether it is measured directly or by FFT, the c value increases (3.64 Å), while a and b decrease (3.64 Å). The difference between c and a or b becomes smaller. At the bottom, we measured 9 atomic planes (Figure 2g). However, in the bottom area (Figure 2g–i), the interplanar spacing of (001) decreases to 3.62 Å while a/b increases to 3.66 Å. Thus, the value of a and b is larger than c. From the middle to the bottom, the change of lattice constant is gradual (Figure S7). At the bottom 15–20 atomic layers, that is 5–6 nm from the surface, there exists vertical compressive strain and lateral tensile strain. Therefore, the range of strain influence decays after leaving the surface. From Figure 2, we can conclude that the vertical strain is compressive strain, and the lateral strain is tensile strain, which is more obvious at the center and near the surface. These results are partially consistent with the previous theoretical studies, that is, tensile stress in the lateral direction and compressive stress in the vertical direction [41, 42]. But the experimental results still have differences compared with the calculation results. The area affected by the lateral tensile stress is much smaller than the reported calculation results. The tensile stress intensity is very limited, and compared with the middle parts, the lattice constant of the center part increases by 0.82%. The vertical compressive stress runs through the entire subunit, which is obvious in the center and bottom (compressed by 0.82%). The difference between the measured results and reported calculation results may come from the migration of twin boundaries and extra atomic planes. Tensile stress is relieved by the migration of twin boundaries and the addition of atomic planes. Solid angle deficiency is filled by elastic strain and extra atomic planes in the (110) facet near the twin boundary. Lateral stress will be quickly released when stress reaches a certain threshold. Consequently, it will not accumulate and appear in a large area. Then (110) facet affects both the lattice constants a and b, and thus the stress on the (100) and (010) facets is relieved simultaneously. As a result, in the center and bottom areas, a = b > c. Therefore, the center and bottom regions can be described as the bct phase, while the middle region is close to the fcc phase, and the entire structure is a bct/fcc heterophase. Similar results can also be obtained in the other subunit (Figure S8). Here, we didn’t use the theoretical value of the Cu lattice constant to compare with the measured values, because different electron microscope equipment may introduce systematic errors when testing. So we only use the tested results of the same sample.

(Figure 3h-k). Based on the secondary electron cut-off energy (Ecutoff ) and equation Wf = hν − (Ecutoff − EF ), the work function (Wf ) of bct/fcc Cu NWs-30, bct/fcc Cu NWs-80, bct/fcc Cu NWs200, and fcc Cu NCs are calculated as 4.58 eV, 4.59 eV, 4.65 eV, and 4,75 eV, respectively (Figure 3l). Here, hν is the energy of laser light used as the incident ultraviolet (UV) photon (21.22 eV), and EF is the Fermi energy [46, 47]. The outside surfaces of Cu NWs and Cu NCs are all (100) facets. Although there are five (111) facets at the ends of Cu NWs, this area is basically negligible in nanowires with extremely high aspect ratios. Thus, the work function change is not caused by different facets. Comparing the 6 of 13

work function of these four samples, electrons are more likely to overflow from the Cu NWs’ surface, which will affect the Cu NWs ability to adsorb intermediates in electrocatalysis. With the diameter increase, the surface work function of Cu NWs increases. Considering the changes in XRD patterns together, this may be because the larger diameter makes the ratio and degree of stress-affected area decrease. The work function reflects the surface state of the material rather than the bulk state. Although higher stress may occur in the center of nanowires with larger diameters, its effect on the surface is likely to be minimal. Figure 3m shows the valence band spectra (VBS) of Cu NWs and Advanced Materials, 2026

FIGURE 3 X-ray and ultraviolet spectral analysis. (a) XRD pattern of bct/fcc Cu NWs-30. (b–e) The zoom-in low-speed scan XRD peaks after peak intensity normalization: (111) peak (b), (200) peak (c), (220) peak (d), and (311) peak (e). Red crystal plane indices indicate the area affected by strain, and the black crystal plane indices indicate the unaffected area. (f) Standard fcc lattice. (g) The bct lattice. (h–k) UPS spectra of bct/fcc Cu NWs-30 (h), bct/fcc Cu NWs-80 (i), bct/fcc Cu NWs-200 (j), and Cu NCs (k). (l) Work functions of bct/fcc Cu NWs and fcc Cu NCs. (m) Valence band spectra (VBS) and d-band center of bct/fcc Cu NWs and fcc Cu NCs. (n) Normalized Cu K-edge XANES spectra of bct/fcc Cu NWs-30, bct/fcc Cu NWs-80, bct/fcc Cu NWs-200, fcc Cu foil, and CuO. (o) Fourier transform of Cu K-edge EXAFS spectra of bct/fcc Cu NWs-30, bct/fcc Cu NWs-80, bct/fcc Cu NWs-200, Cu foil, and CuO. (p) Wavelet transform of Cu K-edge EXAFS spectra for Cu foil (top panel), bct/fcc Cu NWs-30 (middle panel), and bct/fcc Cu NWs-80 (bottom panel).

X-ray absorption spectroscopy (XAS) was applied to further study the electronic structure and local coordination environments of Cu NWs with different diameters. In the Cu K-edge X-ray absorption near-edge structure (XANES) spectra, the white line intensities of Cu NWs are close to those of Cu foil, but far below those of CuO, suggesting that Cu mainly adopts the metallic state along with slight oxidation (Figure 3n). Figure 3o shows the k3 -weighted Fourier-transformed Cu K-edge extended X-ray absorption fine structure (EXAFS) spectra. The dominant peak of Cu foil located at around 2.24 Å is ascribed to the Cu–Cu scattering path. The main peaks of bct/fcc Cu NWs-30, bct/fcc Cu NWs-80, and bct/fcc Cu NWs-200 are basically at the same position, suggesting the dominant metallic state of Cu in three NWs with different diameters. The fitting results show that the Cu-Cu bonds in the Cu NWs with different diameters and Cu foil have a similar bond length of ca. 2.54 Å, but the Cu NWs’ bond length is slightly smaller (Figure S13 and Table S2). The Cu–Cu coordination numbers (C.N.) of bct/fcc Cu NWs-30, bct/fcc Cu NWs-80, bct/fcc Cu NWs-200, and Cu foil are 4.6 ± 0.4, 4.2 ± 0.4, 5.4 ± 0.4, and 12, respectively. This difference might result from the internal defects and the oxidized surface Cu atoms, which reduce the C.N. of Cu–Cu [10, 17, 53]. Simultaneously, the Cu–O coordination numbers in Cu NWs are all much lower than their corresponding Cu–Cu coordination numbers, further proving that the metallic Cu is dominant. The intensity maxima of the wavelet transform (WT) of Cu K-edge EXAFS spectra of Cu NWs are very close to those of Cu foil (Figure 3p; Figure S14).

2.3

Electrocatalytic CO2 RR Performance

The surfaces of both bct/fcc Cu NWs and fcc Cu NCs are almost entirely enclosed by (001) facets. Consequently, there is no facet effect between different samples, and the difference in the electrocatalytic process can be regarded as the difference in the surface electron state, which is due to the bct/fcc heterophase. As a proofof-concept application, the as-prepared bct/fcc Cu NWs and fcc Cu NCs were applied as catalysts in the electrochemical CO2 RR. To better understand the intrinsic performance of Cu materials with different phases, an H-type cell was selected to perform electrocatalytic measurements, as its simplicity can largely avoid errors induced by the complexity of the testing system. The 0.1 M potassium bicarbonate (KHCO3 ) aqueous solution was used as the electrolyte, and the catalysts modified glass carbon electrode (GCE), Ag/AgCl (in saturated KCl solution), and Pt plate were utilized as the working electrode, reference electrode, and counter electrode, respectively. As shown in Figure 4a–d and Figures S15–S18, bct/fcc Cu NWs-30 demonstrate the best performance in CO2 RR. It’s worth noting that bct/fcc Cu NWs-30 deliver the highest FEC2H4 of 56.1% and FEC2+ of 80.2% at -1.15 V (vs RHE). As the diameter increases, the highest FEC2H4 of bct/fcc Cu NWs-80 and bct/fcc Cu NWsAdvanced Materials, 2026

200 doesn’t decrease, which is 55.7% and 54.9% at −1.15 V and −1.10 V (vs RHE), respectively (Figure S19). In contrast, fcc Cu NCs only present the highest FEC2H4 of 48.1%. Meanwhile, bct/fcc Cu NWs-30 nm have the widest C2 H4 generation window, and they can maintain FEC2H4 above 45% in the potential range of −1.05 to −1.30 V (vs RHE). In addition, based on the linear sweep voltammetry (LSV) curves (Figure S20), bct/fcc Cu NWs-30 also show the highest current density compared with the other three samples, demonstrating their excellent electrocatalytic activity. To more clearly show the difference in performance between these samples, the FE of them toward C2 H4 and C2+ at −1.15 V (vs RHE) is compared in Figure 4e. The bct/fcc Cu NWs-30 show the best performance. The FE of C2+ products of bct/fcc Cu NWs-30, bct/fcc Cu NWs-80, bct/fcc Cu NWs-200, and fcc Cu NCs is 80.2%, 77.2%, 64.3%, and 60.3% at −1.15 V (vs RHE), respectively. There is not much difference between bct/fcc Cu NWs-30 and bct/fcc Cu NWs-80 nm. However, bct/fcc CuNWs200 show a decrease in the FE of C2+ products. Although the overall performance of bct/fcc CuNWs-200 nm is still better than that of Cu NCs, the gap is reduced. This phenomenon shows that as the diameter increases, the catalytic ability of the (001) surface decreases significantly. When the diameter is large enough, the performance will be close to a standard fcc lattice. Considering their XRD and UPS results, it is proven that the different abilities of Cu nanowires’ surfaces to generate C2+ products are caused by phase changes inside the material. The greater the internal compressive and tensile strain/stress, the larger the changes induced in the electronic states of surface Cu atoms, which favors the adsorption of intermediates and C–C coupling. When the diameter increases, although the area affected by strain in the center may become larger, the larger diameter gives it more stressrelease areas. As a result, the ratio of the area affected by strain may decrease, and the surface is closer to the standard fcc lattice. Owing to the higher FE and current density, bct/fcc Cu NWs30 deliver a higher partial current density for C2+ products than the other samples (Figure 4f). The C2+ /C1 product ratios were compared to reveal their C-C coupling abilities (Figure 4g). The highest C2+ /C1 ratios of bct/fcc Cu NWs-30, bct/fcc Cu NWs-80, bct/fcc Cu NWs-200, and fcc Cu NCs are 14.9, 11.9, 6.0, and 3.2, respectively, suggesting that unconventional phase Cu NWs with smaller diameters tend to promote C-C coupling. The catalytic stability of bct/fcc Cu NWs-30 was evaluated by long-term CO2 electrolysis in an H-type cell at −1.15 V (vs RHE). After 10 h of operation, the current density displays no obvious decrease, and the FE of C2 H4 remains above 50% (Figure 4h). After the durability test, bct/fcc Cu NWs-30 were characterized by TEM and XRD (Figures S21 and S22). There is little change in the surface. Although some areas show signs of reconstruction, the overall morphology and crystal structure are well preserved, indicating the good catalytic and structural stabilities of bct/fcc Cu NWs. Besides, a systematic comparison with the previously reported catalysts indicates the excellent catalytic performance of bct/fcc Cu NWs toward CO2 RR (Table S3).

2.4

Mechanism Investigation

The electrochemically active surface area (ECSA) was measured by the electrochemical double-layer capacitance (Cdl ) 7 of 13

Cu NCs measured by high-resolution XPS [48–51]. The d-band centers of bct/fcc Cu NWs-30, bct/fcc Cu NWs-80, and bct/fcc Cu NWs-200 are −3.493 eV, −3.416 eV, and −3.184 eV, respectively. They all shift upward compared with fcc Cu NCs (−3.739 eV). The d-band center upshift phenomenon is consistent with the previous study on five-fold twinned structures [52]. This may be caused by the tensile stress on the surface.

(Figure S23). The Cdl of bct/fcc Cu NWs-30 could reach up to 0.131 mF cm−2 , higher than bct/fcc Cu NWs-80 (0.0604 mF cm−2 ), bct/fcc Cu NWs-200 (0.0114 mF cm−2 ), and fcc Cu NCs (0.0660 mF cm−2 ) (Figure 5a). There are two reasons that account for the Cdl decrease with the increase of Cu NWs’ diameter. One is the decrease in specific surface area, and another is the reduction of intrinsic activity. Cu NCs have a higher Cdl value than bct/fcc Cu NWs-80, but their overall performance is much lower than that of bct/fcc Cu NWs-80. Previous studies have shown that there are two types of active sites on the Cu surface. One is more active for the CO2 to CO conversion, while the other favors the further reduction of CO to C2+ products [24]. This indicates that the unconventional bct phase enriches the C-C coupling active sites on the Cu NWs’ surface. As shown in Figure 5b, the CO conversion rate was calculated, bct/fcc Cu NWs-30 and bct/fcc Cu NWs-80 have a similar CO conversion rate of 96.7%, higher than 8 of 13

bct/fcc Cu NWs-200 (95.1%) and Cu NCs (92.1%). Compared with fcc Cu NCs, Cu NWs have a higher CO conversion rate. Figure S24 shows the distribution of C1 products. Compared with fcc Cu NCs, the FE of C1 products decreases significantly on bct/fcc Cu NWs, especially for formic acid. This indicates that bct/fcc Cu NWs tend to suppress formic acid production and improve CO conversion. Considering the reaction pathway (Figure 5c), C1 products usually have two key intermediates that are *OCHO and *COOH, which then further transform to *HCOOH and *CO, respectively [2, 3, 7, 54, 55]. *CO can be further hydrogenated and transformed to *CHO or *COH, which are crucial intermediates for C–C coupling or CH4 [2, 7, 56]. In situ DEMS was applied to probe the dynamic evolution process of electrocatalytic CO2 RR on bct/fcc Cu NWs and fcc Cu NCs (Figure 5d–f). The characteristic mass-to-charge (m/z) ratios for Advanced Materials, 2026

FIGURE 4 Electrocatalytic performance of bct/fcc Cu NWs and fcc Cu NCs. (a–d) FE of CO2 -reduction products obtained on bct/fcc Cu NWs-30 (a), bct/fcc Cu NWs-80 (b), bct/fcc Cu NWs-200 (c), and fcc Cu NCs (d). (e) Comparison of the FEs of C2+ products and C2 H4 generated on bct/fcc Cu NWs and fcc Cu NCs at −1.15 V (vs RHE). (f) The partial current densities of C2+ products on bct/fcc Cu NWs and fcc Cu NCs at different potentials. (g) Comparison of the C2+ /C1 product ratios among bct/fcc Cu NWs and fcc Cu NCs at different potentials. (h) Long-term catalytic stability evaluation of bct/fcc Cu NWs-30 at the constant potential of −1.15 V (vs RHE). The blue arrows indicate the replenishment of electrolyte in the cathode chamber.

the selected fragments are associated with CO (CO, m/z = 28), C2 H4 (C2 H2 + , m/z = 26), and CH4 (CH3 + , m/z = 15) [17, 18, 57–60]. The in situ DEMS patterns of catalysts were collected by linearly sweeping the potential from -0.7 to -1.5 V (vs RHE) with a scan rate of 5 mV/s under CO2 RR conditions. LSV was scanned 5 times, with an interval of 300 s. The data from the three times in the middle were taken. Before the experiment, the electrolyte (0.1 M KHCO3 ) was fully saturated with CO2 . As shown in Figure 5d, the characteristic signals related to CO on bct/fcc Cu NWs-30 appear approximately 40 s later than fcc Cu NCs, and their intensity is much lower than that of the other three samples. Compared with fcc Cu NCs, the peak of bct/fcc Cu NWs is delayed to varying degrees. This indicates that CO on the bct/fcc Cu NWs is more likely to be further rapidly converted into other intermediates. As the ability of fcc Cu NCs to transform CO to other intermediates is limited, CO will stay on their surface for a longer time, and thus makes it easy to be detected. In contrast, the characteristic signals related to C2 H4 show a completely opposite trend (Figure 5e). The peaks of all bct/fcc Cu NWs appear earlier than fcc Cu NCs, suggesting the lower onset potential for C2 H4 generation and the stronger C–C coupling ability. As a competitive product of C2+ products, the appearance of CH4 on all bct/fcc Cu NWs is later than on fcc Cu NCs (Figure 5f). This indicates that the pathway of CH4 is restricted on Cu NWs, and the subsequent reaction Advanced Materials, 2026

pathway of *CO is more conducive to C–C coupling rather than direct reduction to CH4 . The fragment with an m/z ratio of 29 can be assigned to CHO+ (Figure S25), an aldehyde intermediate formed by the continuous hydrogenation of the CO intermediate [61].

2.5

Theoretical Calculations

In order to understand the influences of the presence of the unconventional bct phase of Cu on electrocatalysis, theoretical calculations were also performed to investigate the electronic electroactivity. The surface electronic distributions were first revealed to indicate the bonding and anti-bonding orbitals near the Fermi level (EF ), which are the most electroactive orbitals to undergo electron exchange and transfer during electrocatalysis. For the bct phase Cu surface with strain, there are higher distributions of bonding orbitals, which supply abundant electroactive electrons to promote the CO2 RR toward C2+ products (Figure 6a). In contrast, the fcc Cu surfaces without strain have shown decreased distributions of electroactive bonding orbitals on the surface, lowering the electron transfer efficiency (Figure 6b). To have an in-depth insight into the electronic structures, detailed comparisons of the projected partial density of states (PDOS) 9 of 13

FIGURE 5 Electrochemical mechanism investigation. (a) Fitting results of Cdl for bct/fcc Cu NWs-30, bct/fcc Cu NWs-80, bct/fcc Cu NWs-200, and fcc Cu NCs. (b) CO conversion rates over bct/fcc Cu NWs and fcc Cu NCs at −1.15 V (vs RHE). (c) Schematic diagram of CO2 reduction intermediates. (d–f) In situ DEMS patterns of CO (d), C2 H4 (e), and CH4 (f) in CO2 RR on bct/fcc Cu NWs-30, bct/fcc Cu NWs-80, bct/fcc Cu NWs-200, and fcc Cu NCs. The dashed lines indicate the onset time.

have been supplied (Figure 6c). Based on the two-way crossover linear response method, the evident difference of Cu-3d orbitals is demonstrated, where the bct phase displays evidently upshifted 3d orbitals toward the EF to support the high electroactivity. These results confirm that the strains in Cu NWs with the presence of the bct phase are beneficial to promoting the electroactivity of Cu-3d orbitals. As the diameter of Cu NWs increases, the strain effect becomes less evident, which results in reduced influences on the electronic structures of Cu sites. Figure 6d further shows the site-dependent PDOS of Cu in the strained bct phase. For the bulk Cu sites, the Cu-3d orbitals reveal relatively broad distributions with the co-existence of several low peaks. From the bulk to the surface, the Cu-3d orbitals gradually become more concentrated, which leads to the presence of one high peak with largely increased electron density near the EF . The more concentrated Cu-3d orbitals facilitate the selectivity toward specific 10 of 13

intermediates to promote the CO2 RR. Due to the strain, the bct Cu phase demonstrates a slightly smaller work function of 4.39 eV than that of the fcc phase of 4.69 eV, which is supportive of the experimental characterizations (Figures 6e and 3l). The reduced work function in the bct phase allows easier electron transfers from the Cu surface to the intermediates, promoting the CO2 RR efficiency. For the d-band center, both the surface and bulk of the bct phase exhibit significantly higher d-band centers than those of the fcc phase, representing the enhanced electroactivity induced by the collective contributions of the unconventional bct phase and the strain effect. The LAMMPS simulations demonstrate the atomic strain distributions of Cu NWs with exposed Cu (100) surfaces (Figure 6f), which confirms that the compressive strain exists on the (100) surfaces. Then, the adsorption energies of key intermediates on both bct and fcc Cu surfaces were further compared (Figure 6g). Notably, the bct Cu surface displays lower Advanced Materials, 2026

FIGURE 6 Theoretical study of CO2 RR on bct Cu and fcc Cu. The contour plots of electronic distributions near the Fermi level for (a) bct Cu and (b) fcc Cu. Blue isosurface = bonding orbitals, and green isosurface = anti-bonding orbitals. (c) The PDOS comparisons of Cu-3d orbitals in bct and fcc Cu. (d) The site-dependent PDOS of Cu-3d orbitals in bct Cu. (e) The evolution of d-band centers and work functions in bct and fcc Cu. (f) The atomic strain distributions in fivefold twinned crystal structures of Cu NWs with exposure of (100) surfaces. (g) The adsorption energies of key intermediates during CO2 RR. (h) The reaction trends of different CO2 RR pathways on bct and fcc Cu.

Conclusions

In summary, we have revealed the internal strain distribution and strain release mechanism in fivefold twinned Cu NWs. In the center area and near the surface, each subunit is subjected to vertical compressive strain and lateral tensile strain. In the middle area, the strain basically disappears. The lateral strain is released through the migration of twin boundaries, and the area affected by the strain is limited to a small range. XRD characterizations confirmed the bct/fcc structure of Cu NWs. UPS and XPS tests revealed the different surface electronic states of bct/fcc Cu NWs and fcc Cu. Due to the different degrees of influence of the bct phase, the catalytic performance of Cu NWs is affected to varying degrees. Heterophase bct/fcc Cu NWs-30 show outstanding FEs of 56.2% and 80.2% for C2 H4 and C2+ products at −1.15 V (vs RHE) in an H-type cell, respectively. Nanowires with smaller diameters have a stronger ability to promote C–C coupling during CO2 RR. In situ DEMS patterns verified that, compared to fcc Cu NCs, C2 H4 has lower onset potential on bct/fcc Cu NWs, which indicates that the bct phase is more conducive to C–C coupling. DFT calculations have indicated that the formation of the bct phase with strains has induced evident improvements in the electroactivity for Cu NWs, where the upshifting d-band center, improved binding of key intermediates, and reduced energy barriers all benefit the C2+ product formation during CO2 RR.

Acknowledgements This work was supported by the grant from National Natural Science Foundation of China (Project No. 22175148), grant from Research Grants Council of Hong Kong (Project No. 21309322), grant from Shenzhen Science and Technology Program (Project No. JCYJ20250604184510013), ITC via Hong Kong Branch of National Precious Metals Material Engineering Research Center, grants from City University of Hong Kong (Project No. 9610480, 9610663, 7020103, 7006007, and 9680301), and grant from Guangdong Basic and Applied Basic Research Foundation (Project Advanced Materials, 2026

No. 2024A1515140004). Y.L. acknowledges funding from the research project (Development of advanced catalysts for electrochemical carbon abatement; Project Code: c) is part of the CREATE Thematic Programme in Decarbonisation and is supported by the National Research Foundation, Prime Minister’s Office, Singapore under its Campus for Research Excellence and Technological Enterprise (CREATE) programme.

Conflicts of Interest The authors declare no conflict of interest.

Data Availability Statement The data that support the findings of this study are available from the corresponding author upon reasonable request.

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Supporting Information Additional supporting information can be found online in the Supporting Information section. Supporting File: adma73600-sup-0001-SuppMat.docx.

Advanced Materials, 2026

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