Inch-scale ultrahard diamond wafer with 200 GPa hardness via high-frequency pulsed local non-equilibrium growth

Domain: physics
Published: 2025-12-13
Editorial outcome: major_revision

Abstract
Producing inch-scale, binder-free ultrahard diamond presents a formidable challenge due to the limitations of the conventional high-pressure and high-temperature method. Here, we report a customized microwave plasma chemical vapor deposition technique with high-frequency gas-switching control. By periodically introducing nitrogen, a transient local non-equilibrium growth mode is established, enabling the synthesis of free-standing ultrahard diamond wafers up to 3 mm thick and 5 inches in diameter. The wafers exhibit a Vickers hardness of ~208.3 GPa, comparable to the hardest nano-twinned diamonds, and show exceptional wear resistance—abrasive ratio ~7 times higher than polycrystalline diamond substrate. High-resolution transmission electron microscopy reveals an ultra-dense three-dimensional interlocked stacking fault network (density up to 4.3×10¹² cm⁻²), contributing to superior mechanical properties. This process also allows deposition on commonly used three-dimensional tool surfaces. This work provides a scalable strategy for producing ultrahard, inch-scale diamond suitable for demanding applications in precision machining, semiconductor and aerospace industries.

Introduction
Diamond, as the hardest natural material, is widely used in the field of superhard processing, in the form of polycrystalline diamond (PCD, composite of binder and diamond particles prepared by high temperature and pressure, HPHT), coating, and single crystal diamond (SCD) tools1–7. In the face of rapid advancements across modern industries such as semiconductor8, aerospace9, and mechanical processing10, there is a significantly growing demand for new engineering materials that combine hardness, wear-resistance, and large-size for high-precision machining under extreme conditions. Despite diamond’s remarkable properties, current techniques cannot simultaneously achieve inch-scale, binder-free, ultrahard diamond with hardness surpassing that of natural diamond, particularly in special or complex shapes. This limitation underscores the need for new pathways to develop advanced diamond materials that address the growing demands of modern industries.

The hardness values of typical natural diamond range from 60 to 120 GPa11. In recent years, it possible to produce diamonds that can surpass the hardness of natural ones through nano-structuring strategies. For instance, Irifune et al.12 successfully synthesized ultrahard sintered nano-polycrystalline diamond from graphite under static HPHT with a hardness of 110-140 GPa. Nanocrystalline13,14 and ultra-nanocrystalline diamond (UNCD)15,16 films, characterized by nanoscale grain sizes, typically achieve enhanced hardness through grain boundary strengthening mechanisms. Owing to these mechanical advantages, both NCD and UNCD coatings have been widely employed in cutting and machining tools for their high hardness, excellent wear resistance, and low friction performance17–20. However, the mechanical performance of these films is often constrained by factors such as limited thickness and grain-boundary-related degradation. To overcome these limitations, recent studies have shifted toward atomic-level defect engineering, with nanotwinning emerging as an especially effective strategy. Huang et al.1 successfully synthesized nano-twinned diamonds with an average twin thickness of 5 nm and a hardness of ~200 GPa, which is twice that of natural single crystal diamonds. In the latest study, by controlling the size of onion-like carbon precursors and inducing high-pressure phase transformation, nano-twinned diamonds with an average twin thickness of 2.3 nm were synthesized, achieving a record-breaking hardness of 276 GPa21. However, as a thermodynamic equilibrium method, it demands exceptionally high temperature and pressure, making the synthesis of inch-grade ultrahard diamonds highly challenging. Chemical vapor deposition (CVD), on the other hand, is another widely used method for preparing diamonds by dissociating the gas source to produce plasma, which is usually a localized equilibrium environment during the stable growth of the diamond. In principle, the effective growth space of CVD diamond is much larger, and it offers great flexibility to grow inch-scale diamonds with various features. For instance, via microwave plasma chemical vapor deposition (MP-CVD) of diamond films on silicon wafers, Jing et al.22 recently obtained free-standing, ultrathin flexible diamond membranes with large-area (up to 2-inch) by exfoliation using sticky tape. However, attempts to produce inch-scale diamond wafers with hardness above 200 GPa by the CVD method have so far been unsuccessful.

In this work, we design and develop a customized MP-CVD diamond growth system, as illustrated in Fig. 1. By integrating a fast-response gas-switching control system and a customized-designed chamber into the CVD equipment23, this setup enables precise modulation of the plasma environment during growth, allowing controlled introduction of defects within the diamond lattice. Utilizing this approach, we synthesize large-area diamond wafers on various substrates and investigate their structural and mechanical characteristics. Transmission electron microscopy (TEM) and related techniques reveal extremely high-density stacking fault (SF) structures (defect density of 4.3×10¹² cm⁻²), indicating that high-density stacking-fault diamonds (HSD) are produced by this method. We further explore how the modulation of growth technics governs defect formation and contributes to the mechanical performance of the resulting films. These insights lay the groundwork for scalable fabrication of defect-engineered diamond materials with tailored properties for advanced applications.

Results
Introducing high-density defect structures has been proven to be an effective way for enhancing the hardness of various materials, including diamond1,24–26. In the context of polycrystalline diamond film fabrication, previous studies have shown that incorporating argon or employing methane-rich gas mixtures can significantly influence diamond growth kinetics and defect formation27–30. Among various dopants, nitrogen impurity could be used to reduce grain size31,32 and introduce defects33. Building on these insights, we believe that the key to introducing high density nanostructured defects into CVD-diamond is lies in effectively controlling the nitrogen doping process. To accomplish this, we developed a customized technique for MP-CVD diamond growth: achieving local non-equilibrium growth of diamond through high-frequency cyclic pulsed nitrogen doping, as shown in Fig. 1a. Although cyclic gas-phase doping has been previously applied to modulate electrical properties of diamond, such as in phosphorus doping strategies34–36, our work uniquely applies high-frequency, short-duration nitrogen pulses to induce defects during diamond growth, introducing a new pathway for defect-driven mechanical enhancement. Specifically, by utilizing a high-speed gas valve control system, the nitrogen gas flow switching cycles (with a pulse width of 6 s, see Methods for details) were achieved. Each cycle consists of two phases: the presence and absence of nitrogen gas. These rapid, cyclic fluctuations in gas composition continuously alter the plasma group, thereby creating a local non-thermodynamic equilibrium growth environment (shifting the ratio of CN and C2 radicals periodically, as indicated by the optical emission spectrometer results in Fig. 1b and c), and the periodic variations of growth temperature in Fig. 1d). Using this unique process, we achieved ultra-stable diamond growth on diamond substrates for over 10⁴ cycles. This process ultimately produced a ~3-mm-thick, 5-inch free-standing polycrystalline diamond wafer (Fig. 1e) with controllable thicknesses (see Supplementary Fig. 1, where thickness increased linearly with the number of cycles. The growth rate is 22.1 nm/cycle.). The XRD pattern of the PCD film (Fig. 1f) exhibits a strongest (111) diffraction peak at 44° (predominantly exposed crystal plane), along with weaker (220) and (311) peaks and a very faint (400) peak. The narrow full width at half maximum of the (111) peak suggests high crystallinity. The Raman spectrum (Fig. 1g) shows three characteristic peaks at 1195 cm⁻¹, 1332.5 cm⁻¹, and 1520 cm⁻¹. The sharp peak at 1332.5 cm⁻¹ corresponds to the first-order Raman mode of sp³-bonded diamond. The 1195 cm⁻¹ peak is associated with defect-related or ultrananocrystalline diamond features. The 1520 cm⁻¹ peak may be attributed to nitrogen-induced local lattice distortions, distinct from the typical G band of sp² carbon32,37,38. Surface morphology characterization by Scanning Electron Microscopy (SEM) revealed that most polycrystalline grains exhibit a pyramidal-like surface morphology (Fig. 1h), indicating that each grain predominantly exposed four equivalent {111} planes (consistent with the XRD strongest peak intensity (Fig. 1f) and electron backscatter diffraction (EBSD) result and related discussions (Supplementary Fig. 2). We can precisely adjust the periodical alternating nitrogen doping duration per cycle, and a series of polycrystalline diamond films have been successfully obtained (see Method for specific growth parameters; for ease of discussion, these samples are labeled as HSD-1 to HSD-4).

We then conducted multiple indentation tests (seven times for each sample) on HSD wafers, as well as on conventionally synthesized single-crystalline diamond (SCD) and polycrystalline diamond (PCD). As a case example, Fig. 2a shows the load-hardness curve for the HSD-1 sample under various loads. Vickers hardness tests were conducted under 1 kgf (9.8 N) load to obtain asymptotic hardness39,40, producing indentations with diagonal lengths of approximately 9-10 μm, corresponding to the mesoscopic scale. Multiple indentations were randomly distributed across the sample surface to ensure measurement accuracy and uniformity. Based on this load value, we calculated the Vickers hardness and fracture toughness using standard formulas1,41,42 (see Methods) for all samples, with comparisons shown in Figs. 2b and 2c, respectively. Surprisingly, all HSD wafers exhibited a significant increase in hardness compared to SCD with (100) plane and PCD (Fig. 2b). Furthermore, as the nitrogen-doped diamond growth duration per cycle in step 2 (as described in the Methods section) decreased, the hardness of the HSD wafers increased significantly. Notably, the HSD-1 wafer, which synthesized with the shortest nitrogen doping time per cycle—just 6 seconds—achieved an average Vickers hardness value of 208.3 GPa. The 6-second nitrogen-doped layer duration was optimized through systematic experiments, which showed that shortening the doping time increases hardness without compromising toughness, achieving an optimal balance of mechanical properties. This impressive value is comparable to the reported 5 nm-thick nano-twinned diamond1 and is twice as hard as typical single-crystalline diamond. This meso-scale hardness measurement reliably reflects the macroscopic mechanical performance of the material, especially given the relatively uniform microstructure (Supplementary Fig. 3) and little variation in hardness (180 GPa-210 GPa) observed across the wafer. We also compared the hardness of HSD-1 grown on different substrates (PCD, silicon and SCD) (see Supplementary Fig. 4a) and observed no significant difference in hardness, indicating that this property is intrinsic (not influenced by the substrate). In addition, the comparison of fracture toughness in Fig. 2c and Supplementary Fig. 4b indicates that the increase in hardness does not significantly compromise the toughness of the HSD wafers.

Over the years, the experience of synthesizing ultrahard diamonds has shown us a clear trade-off: the harder the diamond, the more demanding the synthesis conditions, resulting in a more limited (and often smaller) sample size. Here, our analysis of the correlation between hardness values and sample diameters for previously reported ultrahard diamonds (see Fig. 2d) indicates that the size of these diamonds has not exceeded 11 mm, regardless of whether the diamond type is single-crystalline (the hardness range of 70-160 GPa)43–46, polycrystalline (110-140 GPa)12,47–50, nanocrystalline (60-145 GPa)51–54, or nano-twinned with a twin thickness 5 nm (~200 GPa)1. While diamonds with an average twin thickness of 2.3 nm can achieve a hardness of up to 276 GPa21, the sample diameter was only ~1 mm. Encouragingly, the HSD samples in this study overcome previous size limitations, achieving a hardness of 208.3 GPa and a diameter of up to 5 inches.

This periodic nitrogen-doping process cannot only be applied to large-scale planar substrates, but can also be extended to commonly used tools, where uniform diamond coatings can be achieved by optimizing the holder design and plasma exposure conditions. Figure 2e provides a typical example of HSD coating on common engineering tools. After cobalt removal and deposition of an nanocrystalline diamond transition layer, the HSD coating was successfully grown on the surface of the tungsten carbide end mill. SEM images confirm the formation of a uniform and continuous HSD layer across the tool surface. Figure 2f, g exhibit that the HSD sample, when used directly as a cutting tool, can easily cut high-quality single-crystal diamond. Additionally, we compared the wear resistance of various common ultrahard materials used in fabricating cutting tools (see Fig. 2h). The results indicate that the HSD-1 sample, which exhibited the highest hardness, had ~250 times greater wear resistance than polycrystalline c-BN (the most commonly used processing material) and was ~7 times more resistant to wear than polycrystalline diamond (the surface morphology of the samples after wear tests is provided in Supplementary Fig. 5).

The thermal stability of the HSD-1 sample was characterized by thermogravimetry curves measured in air. At a heating rate of 10 °C/min, the onset oxidation temperature was ~760 °C from the thermogravimetry (TG) curve and the differential scanning calorimetry (DSC) curve (Supplementary Fig. 6). The thermal oxidation temperature is comparable to that of natural diamond (770 °C from TG or 720 °C from DSC)1 and greater than that of nanocrystalline diamond (~680 °C)55. This result indicates that the stacking fault structure has excellent thermal stability.

The successful growth of the HSD-1, which exhibits a hardness exceeding 200 GPa, indicates that our unique nano-structuring strategy for producing larger and harder diamonds is effective. To further investigate the origin of the hardness enhancement in the diamond wafer, we conducted an in-depth structural characterization of HSD-1 and other samples using TEM, ranging from nanoscale to atomic scale. The bright-field TEM image in Fig. 3a highlights a key feature within a grain of HSD-1: the presence of extremely high-density, intersecting planar defect structures along two equivalent {111} planes (see the selected area electron diffraction (SAED) inset in the bottom right), which are distinct from previously reported nano-twinned structures1, as they exhibit significantly shorter lengths (average length = 24.7 nm, see Supplementary Fig. 7) and narrower widths. High-resolution TEM (HRTEM) imaging (Fig. 3b) further confirmed that these crystal defects are primarily stacking faults. Specifically, the stacking faults appear in nanometer-sized, bundle-like structures, with each bundle containing numerous individual faults (as indicated by the red arrows in Fig. 3b). These bundles are evenly distributed along two equivalent {111} planes with a high degree of intersection (marked by yellow dashed circles in Fig. 3b), some of which form the well-known Lomer-Cottrell Locks56 (see corresponding atomic-resolution Scanning Transmission Electron Microscopy (STEM) images in Supplementary Fig. 8). We calculated the average stacking fault interval spacing (corresponding to the thickness of nanotwins) along one direction in HSD-1 to be 6.3 nm across multiple areas, as shown in Fig. 3c. And the calculated stacking fault density is 4.3×10¹² cm⁻² (for detailed calculation procedure see Supplementary Information). The shorter nitrogen doping duration per cycle, the higher stacking fault density within the diamond film, as evidenced by the TEM images and corresponding calculated stacking fault densities from HSD-1 to HSD-4 (Supplementary Fig. 9). However, a high-density stacking fault structure can only be achieved when the frequency of periodic nitrogen doping exceeds a certain threshold. This increased stacking fault density directly contributes to higher hardness values, as demonstrated in Fig. 2a and Supplementary Fig. 10. The exceptionally high density of defect structures (corresponding to small stacking fault spacing) and the associated high hardness are consistent with the previously proposed hardening mechanism, which involves the combined contributions of the Hall-Petch and quantum confinement effects57. In addition to typical stacking fault configurations (Supplementary Fig. 11), we also observed several overlap structures (see Supplementary Fig. 12) within the bundles, suggesting that the planar defects are not only short in length and narrow in width but also have very small planar defect areas (averaging 7.3 nm², see Supplementary Fig. 13 for HRTEM views and analysis along the [111] zone axis), resulting in the formation of an ultra-dense, three-dimensional interlocked stacking fault network (Fig. 3d). As shown in Fig. 1h and Supplementary Fig. 2, the majority of the grains in the HSD sample are oriented along the <001> direction. This means that during indentation, the indenter interacts simultaneously with four equivalent {111} planes. And this interlocked network, composed of ultra-dense, intersecting stacking faults uniformly distributed across these four {111} planes (see Fig. 3d), causes dislocations to accumulate at these positions, hindering dislocation movement and effectively preventing deformation, thereby increasing hardness.

This ultra-dense, three-dimensional interlocked stacking fault network has never been reported in diamond materials before. This unique structure is a direct result of the local non-equilibrium growth caused by the disturbance of high-frequency heterogeneous atoms in the plasma, facilitated by our specialized rapidly periodic alternating nitrogen doping cycling technique. The periodic introduction of nitrogen perturbs the plasma and temperature environment, inducing localized non-equilibrium conditions that promote stacking fault formation22. We then further conducted elemental analysis to investigate the distribution of nitrogen in the sample. As shown in the electron energy loss spectroscopy (EELS) results in Fig. 3e, in addition to the diamond C-K edge detected in the perfect-crystalline region (Area 2 in the inset annular dark field (ADF)-STEM image), significant π* and N-K edge signals are also observed in the stacking fault bundle region (Area 1), indicating that nitrogen atoms preferentially accumulate at the stacking faults (also confirmed by energy dispersive x-ray spectroscopy (EDS) results, see Supplementary Fig. 14). Moreover, using atomic-resolution annular bright field (ABF)-STEM (Fig. 3f), we found that the atomic stacking fault layers of HSD-1 exhibit noticeable local tilting, in contrast to the undoped polycrystalline diamond. Therefore, we performed density functional theory (DFT) calculations to determine the formation energies of undoped stacking faults and nitrogen-doped stacking faults in diamond (Fig. 3g). The DFT results reveal that the generalized stacking fault energy (GSFE) of undoped diamond is 5.68 J/m², consistent with literature reports58, whereas nitrogen doping leads to a pronounced reduction to 2.23 J/m². These findings indicate that the local non-equilibrium dynamic growth process, induced by periodically alternating nitrogen doping, continuously perturbs the CVD growth environment of diamond, thereby facilitating the formation of nitrogen-doped stacking faults. The stacking fault formation process can be speculated base on the atom stacking growth mode. In the high-frequency cyclic pulsed local non-equilibrium plasma environment, the periodically doped nitrogen will react with atomic and molecular hydrogen and carbon atoms to form very stable molecules HCN and free radicals CN38,59. The CN radical, acting as a diatomic adsorbent, rapidly forms a stable structure of four atomic islands, which form the nucleus of the subsequent layer of growth on the {111} surface55. As nitrogen accumulates, the lattice becomes distorted. When interstitial nitrogen atoms aggregate into lamellae within the diamond lattice, extrinsic stacking faults are formed through relaxation. Simultaneously, during the preparation process, the dynamic plasma caused by high-frequency periodic nitrogen doping generates energy fluctuations, causing some carbon atoms to overcome the constraints of surrounding atoms and migrate elsewhere. As a result, vacancies form at the original positions. When these vacancies aggregate into plates/disks, intrinsic stacking faults may be formed through collapse relaxation. To sum up, the interstitial nitrogen atoms and local dynamic plasma disturbance in the rapidly periodic nitrogen doping result in a large number of stacking faults.

The nitrogen-doped stacking faults also exhibit relative thermal stability, as evidenced by in situ heating TEM observations (Supplementary Fig. 15), where the thermal degradation of the HSD sample initiates from the perfect-crystalline region rather than the faulted area.

In summary, we have proposed an innovative MP-CVD technique that enables the synthesis of polycrystalline diamond wafers of various thicknesses and sizes (up to 5 inches), while achieving a high hardness exceeding 208.3 GPa. This remarkable enhancement in hardness is primarily attributed to the controllable formation of an extremely high density of stacking faults, caused by the rapidly periodic nitrogen doping strategy. This method significantly reduces the formation energy of stacking faults. Constitutionally, the interstitial nitrogen atoms and local dynamic plasma disturbance cause a large number of stacking faults. This method has led to an unprecedented nanostructured diamond architecture—a highly dense, three-dimensional interlocked stacking fault network—that provides extraordinary resistance to extreme mechanical deformation. This ultrahard CVD diamond technology holds immense potential for a wide range of industrial applications, particularly as cutting tools and coatings for machining the hardest known materials, including diamond, c-BN, and composites containing ultrahard phases. This work provides a new paradigm for enhancing the mechanical performance of diamond, enabling its broader application in aerospace, electronics, and high-precision manufacturing through a defect-engineered strengthening mechanism involving dynamic plasma regulation and stacking fault control.

Methods
Sample preparation
A two-step growth strategy was employed to fabricate large-size, high-thickness free-standing diamond wafers. In the first step, a 5-inch-diameter, 3.5 mm-thick free-standing polycrystalline diamond wafer was deposited on a graphite substrate using a direct current arc plasma jet chemical vapor deposition technique, which offers a high growth rate suitable for thick diamond synthesis. The diamond growth rate ranged from 10 to 15 μm/h23 from the center to the edge of the wafer. The wafer was then mechanically grinded and polished to a final thickness of ~2.8 mm with a smooth surface. In the second step, an ultrahard diamond layer (~200 μm thick) was deposited on the polished wafer using a 915 MHz MP-CVD system with periodically modulated nitrogen doping, under the following growth conditions: microwave power of 11-12 kW, chamber pressure of 11-11.5 kPa, and CH₄/H₂/N₂/O₂ flow rates of approximately 45/900/15/1.5 sccm. The temperature change caused by the gas cycle change is shown in Fig. 1d. During diamond deposition, H₂ maintains a stable plasma and etches the non-diamond phase, while CH₄ provides a carbon source. N₂ doping disrupts the atomic arrangement, while a small amount of O₂ improves the crystal quality by oxidizing the non-diamond carbon. The MP-CVD equipment integrates a fast-response gas flow control system, enabling automatic, periodic adjustments of growth parameters during the diamond deposition process. The programmed gas flow sequence for periodic nitrogen doping is as follows:
(1) Turn on the nitrogen gas flow for 48 seconds.
(2) Turn on the methane gas flow for “x” seconds
(3) Turn off the nitrogen and methane gas flow for 48 seconds.
(4) Turn on the methane gas flow for 48 seconds.
(5) Turn off the methane gas flow for 42 seconds.
These five steps represent one complete cycle for the deposition of HSD films. The durations of step (1) and (3)-(5) were experimentally optimized. The gas flow rates for H₂, CH₄, O₂, and N₂ were set at 300, 15, 0.5, and 5 sccm, respectively. In this work, the mechanical properties of the diamond films were analyzed by varying the nitrogen doping time (“x” in step 2), which was set to 6, 48, 96, and 144 seconds, with the corresponding diamond samples labeled HSD-1 to HSD-4, respectively. Throughout the growth process, optical emission spectroscopy was used to monitor plasma variations, while an infrared thermometer was employed to track temperature changes of the sample.

Characterization of mechanical properties. The hardness and fracture toughness of the prepared diamond films was measured using a Vickers hardness tester (Wilson-WolpertTM Microindentation Tester 401MVD).
The hardness, H, was calculated using the following formula1,41,42:
Hv = 1854.4 × P / a² (1)
where a is the average length of the two diagonals of the indentation, a1 and a2, measured in micrometers (μm), and P is the applied load in newtons (N). The load used in this work was 9.8 N. The unit of the hardness Hv obtained by substituting the values of P and a is GPa.
Fracture toughness was calculated using the following formula1,41,42:
K_Ic = (0.016 ± 0.004) × (E / Hv)^(1/2) × (P / b^(3/2)) (2)
where b = b1 + b2, with b1 and b2 representing the lengths of the two crack diagonals in the indentation, and E is the elastic modulus of the material, with the elastic modulus for diamond samples taken as 1000 GPa1,43. The indentation and crack diagonal lengths were observed and measured using a laser scanning confocal microscopy (OLS4000, Olympus, Japan).
The wear resistance of the samples was evaluated using a fully automated diamond abrasion ratio tester (Changchun Hengyue Electronics Co., Ltd., model HYMH-251). The diamond samples (HSD-1 to HSD-4 grown on PCD substrate) were tilted at 45 degrees and subjected to sliding friction against a diamond composite grinding wheel (13 m/s). The feed rate was set to 0.005 mm every 40 seconds, with the friction load controlled at 300 g. After the test (2.5 hours), the abrasion ratio was determined by the ratio of the abrasion amount of the grinding wheel to the abrasion amount of the samples. A higher abrasion ratio value indicates stronger wear resistance.

DFT calculation. The spin-polarized DFT calculations were performed using the Vienna ab initio simulation package (VASP)60,61 to elucidate the impact of nitrogen doping on the GSFE of the diamond layer. All bulk structures were sourced from the Materials Project database62. The projector augmented wave (PAW) method61, with a cut-off energy of 500 eV, in conjunction with the Perdew-Burke-Ernzerhof (PBE)63 functional, was employed in the DFT calculations. The GSFE is calculated as follows64,65: γ_d = (E_d - E_0) / A, where, E_0 and E_d denote the total energy of the system before and after the slip relative displacement vector d, respectively, and A is the area of the stacking fault. A 112-atom slab model containing an intrinsic stacking fault along the energetically favorable {111} plane was constructed. The model measures approximately 5 × 10 × 13 Å³, with a 15 Å vacuum layer along the z-direction to avoid interactions between periodic images. The topmost and bottom layers were fixed, while the atoms in the other layers were allowed to relax only perpendicular to the slip plane. All models were fully relaxed with an energy convergence criterion of 10⁻⁴ eV and a force convergence criterion of 0.02 eV/Å. The models were constructed using Monkhorst-Pack k-point meshes of (9 × 6 × 1)66.

Microstructural characterization. X-ray diffraction measurements were performed using a Philips X’pert MRD system with a monochromated Cu-Kα X-ray source and a Ge (220) four-crystal monochromator. Raman spectrum was collected using a Renishaw inVia-Qontor micro-Raman imaging system. SEM images were obtained using a field emission scanning electron microscope (Zeiss Gemini 300) equipped with an EBSD detector (Oxford Instruments). TEM samples were prepared using a Thermo Fisher Scientific Helios 5 CX dual-beam system. TEM and STEM images, EDS mappings, and EELS results were acquired using a double Cs-corrected scanning transmission electron microscope (FEI Titan G2 60-300) equipped with a Gatan EELS system and a SuperX EDS system, operating at 300 kV. The in situ heating TEM experiments were carried out using a DENS solutions Wildfire holder with a MEMS-based heating chip.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "The authors present a method using a customized MP-CVD setup with high-frequency pulsed nitrogen doping to create diamond wafers up to 5 inches in diameter and 3 mm thick. They claim a Vickers hardness of ~208.3 GPa, which is comparable to nano-twinned diamonds, showing great scientific significance and application value. The paper’s experimental design is reasonable, the data are detailed, the results are convincing, and the paper as a whole meets the publication criteria of Nature Communications. However, some of the contents need to be further clarified and improved to enhance the rigor and readability of the paper.\n(1) Although the method of “high-frequency pulsed localized non-equilibrium growth” mentioned in the paper is novel, some of the key parameters (e.g., plasma power, temperature fluctuation range, etc.) have not been clearly stated, which may affect the reproducibility of the experiments.\nRecommendations: Add detailed experimental parameters, especially the dynamic regulation mechanism of the plasma environment, and discuss the effects of these parameters on diamond growth and defect formation.\n(2) The manuscript reports seven indentation tests per sample. Could the authors provide data on the spatial uniformity of hardness across the 5-inch wafer (e.g., center vs. edge)? Are there variations in stacking fault (SF) density or nitrogen distribution that might affect mechanical properties at different positions?\n(3) The gas-switching cycle (6s nitrogen pulse) is critical. Was this parameter optimized systematically, or is it based on empirical observation? Is there a theoretical framework guiding the choice of pulse frequency?\n(4) What is the thermal stability of stacked layer error networks? Does defect annihilation occur in high temperature applications?\n(5) The TEM analysis emphasizes SFs, but could other defects (e.g., dislocations, twins) contribute to hardening? Are the SAED streaks in Fig. 3a exclusively attributable to SFs, or might they indicate additional planar defects?\n(6) The DFT calculations show reduced formation energy for N-doped SFs. What slab/supercell configurations were used? Were multiple SF orientations or nitrogen positions tested to ensure the generality of the conclusion?\n(7) Is the economic cost of this method competitive with diamonds prepared by the conventional HPHT method?", "reviewer_2": "The manuscript by J. Tu et al. presents a study on the influence of high-frequency gas-switching, particularly nitrogen inlet modulation, on the growth kinetics of diamond thin films. The authors report the fabrication of ultrahard diamond wafers at the inch scale, which is technologically significant. However, several concerns and points of clarification must be addressed before publication. I recommend acceptance after major revision, contingent upon the incorporation of the following experimental details and clarifications.\nComments:\nSome sentences in the abstract and introductory part of the manuscript are repeated. It is proposed that the appropriate text parts be reformulated.\nReferences on the growth of nanocrystalline and ultrananocrystalline diamond (UNCD) films should be included. These materials are conceptually similar to the ones discussed in the manuscript and generally show comparable hardness. Including these references would better contextualize the findings.\nThe use of nanocrystalline diamond films in cutting tools is already well-established. A critical review of existing applications and comparative analysis should be added to enhance the relevance of the current results.\nPlease verify if “3 mm” refers to the actual thickness of the grown diamond. This is unusually high for CVD-grown films and requires confirmation or correction, while thinner samples were later tested.\nThe statement “can be additionally grown on surfaces of other complicated engineering tools” is unclear. In microwave plasma systems, the plasma is typically localized in a spherical or disc shape, which limits uniform coating on complex 3D geometries, especially along the Z-axis. Please clarify or revise this statement.\nThe manuscript presents nitrogen as the sole effective parameter for stacked-diamond formation. Please comment on the potential role of other gases, such as argon-based plasmas or methane-oversaturated gas mixtures with/without nitrogen, which are also reported to affect the growth mechanisms in a similar manner.\nProvide the growth rate corresponding to the reported 3 mm thick diamond layer. This would substantiate claims regarding the high growth efficiency.\nA more detailed discussion of the XRD peak intensities and the corresponding crystallographic information is needed.\nSimilar detailed analysis should be applied to other techniques used (e.g., Raman spectroscopy). Spectral features and peak assignments should be scientifically interpreted rather than descriptively reported.\nThe comparison of hardness should be presented with better clarity: a) Specify the substrate sizes, especially for SCD samples and b) since hardness tests are performed at the micro-scale, additional commentary on macro-scale hardness implications should be included and finally, c) it is not clear the presented hardness values of PCD and SCD substrates were measured by the authors using the same testing methods as those used for the current samples (please, clarify or measure PCD/SCD in your setup), while reliance on external literature values may lead to inconsistent or misleading comparisons.\nThe growth of diamond films on 5-inch planar substrates using microwave plasma CVD at 915 MHz is not currently a technological breakthrough. Phrases such as “5 inch impressive” are considered overstated and should be omitted or rephrased.\nThe information regarding direct HSD growth on WC inserts is potentially misleading due to the catalytic graphitization effect of cobalt in WC. Either elaborate on any barrier layers or initial pretreatments of WC used, or revise this statement accordingly.\nToughness values calculated using elastic modulus may be overestimated. Nanocrystalline diamond materials can exhibit values 10–25% lower than theoretical estimations. Please revise and discuss the limitations of this method.\nSome labels in Figure 2 are not clearly linked to specific experimental results or references. Please revise for better clarity and traceability.\nThe concept of cyclic gas-phase doping has been previously explored, particularly for phosphorus doping. Relevant studies should be cited, and the novelty of the current doping method should be more clearly positioned.", "reviewer_3": "This manuscript presents a novel approach for synthesizing large-scale ultrahard diamond wafers using a customized microwave plasma chemical vapor deposition (MP-CVD) method with high-frequency pulsed local non-equilibrium growth. The authors report achieving diamond wafers up to 5 inches in diameter with exceptional hardness (~208.3 GPa) through periodic nitrogen doping, which they attribute to the formation of a three-dimensional interlocked stacking fault network. This manuscript presents potentially significant advances in ultrahard diamond synthesis with impressive scale and claimed properties. However, several critical aspects require further clarification and validation before publication. I recommend major revision to address the concerns outlined in the following, particularly regarding the hardness measurement methodology, stacking fault characterization, and stacking fault energy calculations. If these concerns can be adequately addressed, this work would represent a valuable contribution to the field of ultrahard materials and diamond synthesis.\n1. The reliability of the hardness measurements is questionable based on the evidence presented. The Vickers indentation shown in Fig. 2a inset displays obvious features with ridges extending beyond the marked area, suggesting potential overestimated hardness. The reported PCD hardness (~100 GPa) is significantly higher than typical values (~70 GPa), raising questions about the calibration of the measurement system. Additionally, the authors don't specify which crystallographic plane of the single-crystal diamond was measured, which is critical since diamond hardness varies significantly with crystallographic orientation. The reported fracture toughness for SCD (~9.1 MPa·m^(1/2)) is also unusually high compared to typical values (less than 5 MPa·m^(1/2)), further suggesting potential measurement inconsistencies.\n2. The stacking fault characterization presented in the manuscript requires substantial improvement. Many of the TEM images lack sufficient resolution and clarity to unambiguously identify stacking fault structures. The quasi-periodic structures visible in Fig. 3b warrant higher-resolution atomic-scale characterization to confirm their nature. While the authors report a stacking fault density (usually defined as SF area/volume) of 4.3×10^12 cm^-2, the statistical sampling methodology is not clearly described, raising questions about the representativeness of this value. Given that the stacking fault network is presented as the primary mechanism for the enhanced hardness, more comprehensive and higher quality microstructural characterization is essential.\n3. There is a concerning ambiguity between the main text and the Methods section regarding which parameter is being varied to control nitrogen doping. The main text states \"as the nitrogen doping duration per cycle decreased,\" but the Methods section indicates that the N₂ turn-on duration remained constant at 48 seconds. The variable parameter appears to be the methane gas flow duration after nitrogen introduction (parameter \"x\" in step 2), which should be clarified to ensure reproducibility of the process.\n4. The stacking fault energy calculations presented in the manuscript require further validation. The calculated stacking fault energy values appear to be in the range of hundreds of mJ/m², which is much lower than typical values for diamond (~10 J/m²). Complete generalized stacking fault energy (GSFE) calculation results should be provided as supplementary information to validate these values. Furthermore, the thermal stability of these high-density stacking fault structures is not addressed. Heat treatment experiments would be valuable to understand how this microstructure might evolve under thermal loading, which is crucial for assessing the practical utility of these materials.\n5. The manuscript would benefit from additional language editing to improve clarity and readability. There are several instances of grammatical errors and awkward phrasing throughout the text."}, "author_response": "Responses to Reviewers’ Comments\nInch-Scale Ultrahard Diamond Wafer via High-Frequency\nPulsed Local Non-Equilibrium Growth\nREVIEWER COMMENTS\nReviewer #1 (Remarks to the Author):\nThe authors present a method using a customized MP-CVD setup with high-frequency\npulsed nitrogen doping to create diamond wafers up to 5 inches in diameter and 3 mm\nthick. They claim a Vickers hardness of ~208.3 GPa, which is comparable to nano-\ntwinned diamonds, showing great scientific significance and application value. The\npaper’s experimental design is reasonable, the data are detailed, the results are\nconvincing, and the paper as a whole meets the publication criteria of Nature\nCommunications. However, some of the contents need to be further clarified and\nimproved to enhance the rigor and readability of the paper.\nResponse: We sincerely appreciate the reviewer for his/her time and effort on\nreviewing our manuscript and offering the positive feedback and constructive\ncomments on our work.\nWe have now carefully addressed your carefully addressed your comments/suggestions and incorporated\nthe corresponding revisions in the manuscript. The text changes are highlighted in\nyellow for your convenience. Please also find our point-by-point responses below.\n1. Although the method of “high-frequency pulsed localized non-equilibrium growth”\nmentioned in the paper is novel, some of the key parameters (e.g., plasma power,\ntemperature fluctuation range, etc.) have not been clearly stated, which may affect the\nreproducibility of the experiments.\nRecommendations: Add detailed experimental parameters, especially the dynamic\nregulation mechanism of the plasma environment, and discuss the effects of these\nparameters on diamond growth and defect formation.\nResponse 1: Thank you for your valuable suggestions. In the revised manuscript, we\nhave added the detailed experimental parameters in the Methods section to enhance\nreproducibility. These include:\nMicrowave plasma power: 11-12 kW\nChamber pressure: 11-11.5 kPa\nGas flow rates (CH₄/H₂/N₂/O₂): ~45/900/15/1.5 sccm\nThe growth temperature fluctuates periodically in the range of 775-830℃\naccording to the changes in the gas during the cycle (Fig.1d).\nWe have also elaborated on the role of each gas and the dynamic regulation\nmechanism of the plasma environment in the revised Discussion section. Specifically,\nH₂ serves as the main carrier gas, ensuring stable plasma discharge and etching non-\ndiamond phases via H radicals to retain sp³-bonded diamond. CH₄ provides the carbon\nsource for diamond growth. N₂ acts as a doping gas to disrupt the atomic stacking\nsequence and induce stacking faults. A small amount of O₂ helps oxidize non-diamond\ncarbon, improving crystalline quality. The pulsed introduction of N₂ creates periodic\nperturbations in the plasma composition and substrate temperature, resulting in\nlocalized non-equilibrium growth conditions. These fluctuations promote the formation\nof stacking faults and other defects, thereby enhancing the mechanical properties of the\ndiamond films.\nWe believe these revisions have clarified the experimental setup and mechanism,\nand we appreciate your suggestion, which has led to a more rigorous and transparent\npresentation of our method.\nWe have added relevant content in the Methods and Discussion sections in the\nrevised manuscript as follows.\n“During diamond deposition, H₂ maintains a stable plasma and etches the non-\ndiamond phase, while CH₄ provides a carbon source. N₂ doping disrupts the atomic\narrangement, while a small amount of O₂ improves the crystal quality by oxidizing the\nnon-diamond carbon.” (Please see Page 18 in the revised manuscript)\n“The periodic introduction of nitrogen perturbs the plasma and temperature\nenvironment, inducing localized non-equilibrium conditions that promote stacking fault\nformation.” (Please see Page 13 in the revised manuscript)\n“An ultrahard diamond layer (~200 μm thick) was deposited on the polished wafer\nusing a 915 MHz MP-CVD system with periodically modulated nitrogen doping, under\nthe following growth conditions: microwave power of 11-12 kW, chamber pressure of\n11-11.5 kPa, and CH₄/H₂/N₂/O₂ flow rates of approximately 45/900/15/1.5 sccm. The\ntemperature change caused by the gas cycle change is shown in Fig.1d.” (Please see\nPage 18 in the revised manuscript)\n2. The manuscript reports seven indentation tests per sample. Could the authors provide\ndata on the spatial uniformity of hardness across the 5-inch wafer (e.g., center vs. edge)?\nAre there variations in stacking fault (SF) density or nitrogen distribution that might\naffect mechanical properties at different positions?\nResponse 2: In our initial tests, the hardness values were obtained from randomly\nselected positions across the sample surface. To further evaluate the spatial uniformity\nof hardness across the 5-inch wafer, we performed additional Vickers indentation tests\nat 12 evenly distributed positions from the center to the edge, as shown in Figure R1.\nThe corresponding results are summarized in Table R1. These measurements reveal a\nslight increase in hardness from the center toward the edge of the wafer; however, the\noverall variation remains within the margin of error, indicating good spatial uniformity.\nFigure R1 12 indentations evenly spaced from center to edge on a 5-inch diamond\nwafer\nTable R1 The diagonal length, crack length, hardness and toughness of 12 indentations\nPositions a 1 /μm a 2 /μm b 1 /μm b 2 /μm Hv/GPa K Ic /MPa·m1/2\n1 10.063 9.912 37.938 32.616 182.2 5.0\n2 10.089 9.721 29.972 33.887 185.2 5.7\n3 9.672 9.357 29.777 36.038 200.8 5.2\n4 9.962 10.201 33.649 26.181 178.8 6.4\n5 9.769 9.656 38.33 37.288 192.6 4.3\n6 9.321 9.371 35.894 33.886 208.1 4.7\n7 10.211 10.199 37.632 36.188 174.5 4.7\n8 9.877 9.841 38.363 37.572 187.0 4.4\n9 9.336 9.54 35.385 36.725 204.0 4.5\n10 9.992 10.007 36.534 36.357 181.7 4.7\n11 9.658 9.66 30.257 35.489 194.8 5.3\n12 9.34 9.24 34.245 35.575 210.6 4.7\nThe stacking fault density can influence the mechanical properties. In fact, we also\nperformed TEM characterizations at both central and edge regions (Supplementary Fig.\n3). High-density stacking faults were observed in both locations. Since nitrogen doping\npromotes the formation of stacking faults, its distribution can influence mechanical\nproperties. However, the nitrogen supply during the process was sufficient and well-\ndistributed, preventing significant variations across the wafer. To improve the clarity\nand transparency of the hardness uniformity and its potential reasons, we have added\nthe structure of the center and edge in the revised manuscript (Supplementary Fig. 3).\n“Supplementary Fig. 3 Microstructural analysis of the 5-inch ultra-hard diamond\nfilm at different positions. a, b HRTEM images of sample taken from the center region\nof the 5-inch film. c, d HRTEM images of sample taken from the edge region. Both the\ncenter and edge regions exhibit the same stacking fault structures and similar stacking\nfault densities.” (Please see Page 6 in the revised Supplementary Information)\n3. The gas-switching cycle (6s nitrogen pulse) is critical. Was this parameter optimized\nsystematically, or is it based on empirical observation? Is there a theoretical framework\nguiding the choice of pulse frequency?\nResponse 3: We appreciate your insightful comment regarding the gas-switching cycle\noptimization. The 6-second nitrogen pulse duration was indeed determined through\nsystematic experimental optimization rather than empirical observation alone. In fact,\nthe choice of pulse frequency in our study is guided by both experimental results and\nan underlying theoretical framework related to defect modulation in multilayer\ndiamond structures.\nPreviously, in our work on single-crystal diamond 1, we found that periodically\nnitrogen-doped multilayer structures could significantly enhance fracture toughness.\nMotivated by these results, and considering that polycrystalline diamond typically\nexhibits lower toughness than single-crystal diamond, we aimed to apply the same\nstrategy to polycrystalline diamond to improve its mechanical properties. During this\nprocess, we systematically investigated multilayer diamond films composed of\nalternating nitrogen-doped and undoped layers with equal thickness. By varying the\nperiod (i.e., individual layer thickness), we observed that while the fracture toughness\nimproved slightly, the hardness was more sensitive to the layer thickness and increased\nsignificantly as the period decreased. However, since the nitrogen-doped layers tend to\nform ultra-nanocrystalline structures containing amorphous carbon, which could limit\nhardness improvement, we further optimized the doping cycle by reducing the growth\ntime of the nitrogen-doped layer while keeping the undoped layer constant. Through\nthese experiments, we found that when the nitrogen-doped layer growth time was\nreduced to 6 seconds, the resulting films exhibited the highest hardness without\nsacrificing toughness. This time point was selected as the optimal nitrogen pulse\nduration. Of course, in addition to the growth time of the nitrogen-doped diamond layer,\nother process parameters were also experimentally optimized. We appreciate your\nsuggestions, which allowed us to further clarify the rationale and optimization pathway\nbehind our nitrogen pulsing strategy. To provide further clarity on this issue, the above\nanalysis has been included in the Supplementary Text (Please see Page 2 in the revised\nSupplementary Information). We have added relevant content in the revised\nmanuscript as follows.\n“The 6-second nitrogen-doped layer duration was optimized through systematic\nexperiments, which showed that shortening the doping time increases hardness without\ncompromising toughness, achieving an optimal balance of mechanical properties.”\n(Please see Page 8 in the revised manuscript)\n1. Zhao Y, et al. Toughness enhancement of single-crystal diamond by the\nhomoepitaxial growth of periodic nitrogen-doped nano-multilayers. International\nJournal of Minerals, Metallurgy and Materials, 30(4), 766-771 (2023).\n4. What is the thermal stability of stacked layer error networks? Does defect\nannihilation occur in high temperature applications?\nResponse 4: Following your suggestion, we evaluated the thermal stability of the ultra-\nhard polycrystalline diamond film with a high-density stacking fault structure in air\nusing a thermogravimetric analyzer (TGA55) at a heating rate of 10°C/min. The results\nare shown in the newly added Supplementary Fig. 6. The thermogravimetry (TG) curve\nshows that the initial oxidation temperature of diamond is 755°C, and the differential\nscanning calorimetry (DSC) curve shows that the initial oxidation temperature of\ndiamond is 760 °C. The thermal oxidation temperature is comparable to that of natural\ndiamond (770 ℃ from TG or 720 ℃ from DSC) and greater than that of nanocrystalline\ndiamond (~680 °C).\nIn addition, to address your concerns regarding the annihilation behavior of defect\nstructures at high temperatures, we performed detailed in situ heating TEM experiments\non the HSD sample at atomic resolution, which specifically focused on the structural\nevolution of stacking faults at elevated temperatures (from room temperature to\n1100 °C). As shown in Supplementary Fig. 7, the stacking fault structures exhibit\nexcellent thermal stability, remaining intact up to 900 °C. Notably, degradation of the\nHSD sample begins from the perfect crystalline regions (yellow dashed lines), rather\nthan from the stacking fault regions (indicated by red arrows). The annihilation of\ndefects is observed at 1000 °C in the TEM. The in situ results clearly demonstrate that\nthe thermal stability of HSD is not compromised by the stacking fault structures.\nThese results indicate that the stacking fault structure has excellent thermal\nstability. In the revised version, we have added the data and discussions as following:\n“The thermal stability of the HSD-1 sample was characterized by thermogravimetry\ncurves measured in air. At a heat ingrate of 10 ℃/min, the onset oxidation temperature\nwas ~760 ℃ from the thermogravimetry (TG) curve and the differential scanning\ncalorimetry (DSC) curve (Supplementary Fig. 6). The thermal oxidation temperature is\ncomparable to that of natural diamond (770 °C from TG or 720 °C from DSC)1 and\ngreater than that of nanocrystalline diamond (~680 °C)56. This result indicates that the\nstacking fault structure has excellent thermal stability.” (Please see Page 10 in the\nrevised manuscript)\n“Supplementary Fig. 6 Typical thermal stability of HSD-1 sample. TG a and DSC b\ncurves were measured in air at a heating rate of 10 ℃/min.” (Please see Page 9 in the\nrevised Supplementary Information)\n“1. Huang Q., Yu D., Xu B., Hu W., Ma Y., Wang Y., et al. Nanotwinned diamond with\nunprecedented hardness and stability. Nature, 510 (7504), 250-253 (2014).\n56. Solozhenko, V. L., Kurakevych, O. O. & Le Godec, Y. Creation of nanostuctures by\nextreme conditions: high-pressure synthesis of ultrahard nanocrystalline cubic boron\nnitride. Adv. Mater. 24, 1540–1544 (2012).” (Please see Page 22 and 24 in the revised\nmanuscript)\n“In situ heating TEM observations (Supplementary Fig. 7) reveal that the presence of\nstacking faults does not compromise the thermal stability of the HSD diamond, as lattice\ndegradation initiates from the perfect crystalline regions rather than the faulted areas.”\n(Please see Page 10 in the revised manuscript)\n“Supplementary Fig. 7 In situ heating TEM investigation of the thermal stability of\nstacking faults in HSD. a-l HRTEM image sequence showing the evolution of stacking\nfaults in HSD at elevated temperatures (from room temperature up to 1100 °C). The\ntemperature was increased in steps of 100 °C, with a holding time of 10 min at each\nstep. Stacking fault regions are indicated by red arrows, and the degraded regions are\noutlined with yellow dashed lines. The onset temperature of degradation is slightly\nhigher than that observed in TG and DSC measurements on bulk samples, which may\nbe attributed to the vacuum environment within the TEM.” (Please see Page 10 in the\nrevised Supplementary Information)\n5. The TEM analysis emphasizes SFs, but could other defects (e.g., dislocations, twins)\ncontribute to hardening? Are the SAED streaks in Fig. 3a exclusively attributable to\nSFs, or might they indicate additional planar defects?\nResponse 5: With regard to other defects, we note that perfect dislocations were\nscarcely observed in our samples. Partial dislocations were consistently found at the\nterminations of stacking faults, as shown in Supplementary Fig. 12. Although a few\ntwin structures were also identified, their density was considerably lower than that of\nthe nitrogen-doped stacking faults, which are a characteristic feature of our samples.\nWhile it is well known that twins can contribute to hardening 1, 2, in our case, the high-\ndensity stacking faults play the dominant role in enhancing hardness.\nFor the SAED streaks in Fig. 3a, we believe that they are predominantly\nattributable to SFs, which is consistent with numerous previous studies reporting the\nappearance of streaks in SAED as a typical signature of SFs 3, 4. In contrast, other types\nof planar defects, such as twins, would be expected to produce additional, discrete\ndiffraction spots rather than streaks in the SAED pattern 1, 5.\n1. Huang Q, et al. Nanotwinned diamond with unprecedented hardness and stability.\nNature, 510(7504), 250-253 (2014).\n2. Ying P, et al. Enhancing the hardness of diamond through twin refinement and\ninterlocked twins. Nature Synthesis, 4(3), 391-398 (2025).\n3. Bae H J, et al. On the stacking fault forming probability and stacking fault energy\nin carbon-doped 17 at% Mn steels via transmission electron microscopy and atom\nprobe tomography. Journal of Materials Science & Technology, 115, 177-188\n(2022).\n4. Krishnan B, et al. Substrate-Dependent Orientation and Polytype Control in SiC\nNanowires Grown on 4H-SiC Substrates. Crystal growth & design, 11(2), 538-541\n(2011).\n5. Zhu, W, et al. Structural imperfections in CVD diamond films. Journal of Materials\nResearch, 1989, 4: 659-663.\n6. The DFT calculations show reduced formation energy for N-doped SFs. What\nslab/supercell configurations were used? Were multiple SF orientations or nitrogen\npositions tested to ensure the generality of the conclusion?\nResponse 6: In our DFT calculations, we employed slab models containing intrinsic\nstacking fault (SF) structures based on {111} planes, which are the most energetically\nfavorable and commonly observed in diamond. The model has dimensions of\napproximately 5 × 10 × 13 Å3, with a 15 Å vacuum layer introduced along the out-of-\nplane (z) direction to eliminate spurious interactions between periodic images. The slab\nconsists of 112 atoms in total. The Methods section of the manuscript already includes\nsome descriptions of the models. To improve clarity for readers, we have added the\nfollowing supplementary details in the revised manuscript.\n“A 112-atom slab model containing an intrinsic stacking fault along the energetically\nfavorable (111) plane was constructed. The model measures approximately 5 × 10 ×\n13 Å3, with a 15 Å vacuum layer along the z-direction to avoid interactions between\nperiodic images.” (Please see Page 20 in the revised manuscript)\nDuring diamond growth, atomic stacking primarily occurs along the close-packed\n(111) planes. Periodic nitrogen doping tends to disrupt this stacking sequence, leading\nto the formation of stacking faults along these planes. In contrast, stacking faults are\nunlikely to form on other crystallographic orientations due to their unfavorable stacking\nconfigurations. Therefore, our calculations focused exclusively on the stacking fault\nenergy along the (111) planes.\nTo evaluate the effect of nitrogen doping positions, we systematically substituted\nnitrogen atoms at multiple plausible sites within the stacking fault plane, as shown in\nSupplementary Fig. 16. Each configuration was fully relaxed to obtain its formation\nenergy. The stacking fault energies for the three different nitrogen positions are 0.129\nJ/m2, 0.128 J/m2, and 0.129 J/m2, respectively. These results indicate that the specific\nposition of the nitrogen atom has little effect on the stacking fault energy, but its\nincorporation consistently reduces the stacking fault energy compared to the undoped\ncase. To be more rigorous, we have added relevant content in the revised manuscript as\nfollows.\n“In addition, three representative nitrogen doping configurations within the stacking\nfault plane were evaluated (Supplementary Fig. 16), yielding stacking fault energies of\n0.129, 0.128, and 0.129 J/m2. The results show that nitrogen position has little effect,\nbut its incorporation consistently reduces the stacking fault energy compared to the\nundoped case.” (Please see Page 14 in the revised manuscript)\n“Supplementary Fig. 16 Model of stacking fault structure with different nitrogen\npositions in diamond. Blue and pink spheres represent carbon and nitrogen atoms,\nrespectively.” (Please see Page 19 in the revised Supplementary Information)\n7. Is the economic cost of this method competitive with diamonds prepared by the\nconventional HPHT method?\nResponse 7: Both MP-CVD and HPHT are common methods for preparing diamonds.\nTheir principles for preparing diamonds are different, and the obtained diamonds have\ndifferent sizes and qualities, resulting in different application fields. In terms of\neconomic costs, while the HPHT method is widely used for producing bulk diamonds,\nit typically involves higher per-experiment costs, stringent equipment and ultrahigh\npressure environmental requirements, and expensive high-pressure components (e.g.,\nanvils, reaction cells). These factors significantly increase the overall cost, especially\nfor producing high-quality single crystals. In contrast, the MP-CVD method used in our\nstudy offers clear cost advantages in several aspects. First, it operates under relatively\nmoderate temperature and pressure conditions, which reduces the requirements on\nequipment and laboratory infrastructure. Second, MPCVD systems support batch\ngrowth, flexible size control, and reduced material waste, allowing for more efficient\nuse of source gases and longer operational lifetimes of key components. Moreover, the\nlower maintenance costs, ease of automation, and potential for continuous or scalable\nproduction make MPCVD increasingly attractive from an economic standpoint,\nparticularly for applications demanding high crystal quality, precise doping (e.g., B or\nN), and defect engineering. These characteristics make MPCVD-prepared diamonds\nmore suitable for high-end applications (e.g., in electronics, optics, and precision\ncutting tools). Therefore, considering both quality and cost-efficiency, the MPCVD\nmethod demonstrates strong competitiveness, and in many cases, a clear economic\nadvantage over the conventional HPHT technique, especially for research-grade or\ndevice-oriented diamond materials.\nBased on the traditional MP-CVD, combined with the pulsed non-equilibrium plasma\ngrowth technology, we introduced a high-density three-dimensional interlocking\nstacking fault structure in diamond, which can produce ultra-high hardness and large-\nsized diamonds. It can be used for ultra-precision cutting tools, special mechanical parts,\netc. For example, in the fields of aerospace and national defense, large-sized diamonds\ncan be used to manufacture key components such as wear-resistant bearings and high-\nprecision pistons, taking into account both lightweight and high strength. In addition,\nlarge-sized diamonds can be cut into many small-sized diamonds and special-shaped\ntools through laser cutting, thereby reducing costs.\nLastly, we thank you once again for your valuable comments and suggestions that\nreally help improve our manuscript.\nReviewer #2 (Remarks to the Author):\nThe manuscript by J. Tu et al. presents a study on the influence of high-frequency gas-\nswitching, particularly nitrogen inlet modulation, on the growth kinetics of diamond\nthin films. The authors report the fabrication of ultrahard diamond wafers at the inch\nscale, which is technologically significant. However, several concerns and points of\nclarification must be addressed before publication. I recommend acceptance after major\nrevision, contingent upon the incorporation of the following experimental details and\nclarifications.\nResponse: We firstly thank the reviewer for your overall positive feedback and\nthoughtful evaluation of our manuscript with many critical and constructive suggestions.\nBelow, we provide a point-by-point response to the comments/concerns raised. All\nrevisions in the manuscript are highlighted in yellow for your further review.\nComments:\n1. Some sentences in the abstract and introductory part of the manuscript are repeated.\nIt is proposed that the appropriate text parts be reformulated.\nResponse 1: As you mentioned, some sentences in the Abstract and Introduction were\nindeed repetitive, which may have affected the clarity and conciseness of the\nmanuscript. In the revised manuscript, we have carefully rephrased and streamlined\nthese sentences to eliminate unnecessary repetition while preserving the clarity and\ncoherence of the main content.\nThese changes can be found in the revised manuscript. (Please see Page 2 and Page 4\nin the revised manuscript)\nWe believe these modifications have significantly improved the readability of the\nmanuscript, and thank you for highlighting this issue.\n2. References on the growth of nanocrystalline and ultrananocrystalline diamond\n(UNCD) films should be included. These materials are conceptually similar to the ones\ndiscussed in the manuscript and generally show comparable hardness. Including these\nreferences would better contextualize the findings.\nResponse 2: We thank you very much for this suggestion regarding the inclusion of\nreferences on nanocrystalline (NCD) and ultrananocrystalline diamond (UNCD) films.\nWe agree that NCD and UNCD films are relevant to the context of our work. In the\nrevised manuscript, we have included relevant references on the synthesis and\nmechanical performance of NCD and UNCD films. NCD and UNCD films are typically\nsynthesized via CVD methods and are characterized by grain sizes in the nanometer or\nsub-10 nm range, with mechanical strengthening mainly arising from grain boundary\nhardening. Representative works include 1, 2 demonstrated UNCD with exceptional\nhardness and tribological properties, and Wang et al. and Lin et al. 3, 4, who reported\nscalable growth of hard nanocrystalline diamond coatings. Building upon this grain-\nboundary-strengthening strategy, more recent research has advanced into defect\nengineering at the atomic scale, particularly through nanotwinning. For instance, the\nnano-twinned diamond (NTD) reported by Tian et al. 5, 6 achieves ultrahigh hardness\n(>200 GPa) by introducing dense coherent twin boundaries that serve as effective\nbarriers to dislocation motion. These studies highlight how structural refinement from\nthe grain to the twin level can drastically enhance diamond’s mechanical performance.\nBuilding on these insights, our work introduces a novel approach to tailoring the\nstacking fault structures in diamond films through periodically modulated nitrogen\ndoping during MP-CVD growth. Our work extends the concept of nanostructured\ndiamond strengthening by introducing a new pathway—defect-engineered hardness\nthrough dynamic plasma regulation and stacking fault engineering—which is\nfundamentally different from traditional grain refinement in NCD and UNCD, and may\nopen new directions for diamond film design.\n1. Gruen, D. M. Nanocrystalline diamond films. Annual Review of Materials Science\n29, 211-259 (1999).\n2. Sumiya H., Irifune T. Microstructure and mechanical properties of high-hardness\nnano-polycrystalline diamonds. SEI Technical Review-English Edition, 66, 85\n(2008).\n3. Lin C. R., Liao W. H., Wei D. H., et al. Formation of ultrananocrystalline diamond\nfilms with nitrogen addition. Diamond and Related Materials, 20 (3), 380-384\n(2011).\n4. Tu J, Zhao Y, Chan S, et al. Microstructure control and mechanical properties of\nultra-nanocrystalline diamond films. Ceramics International, 50 (21), 43780-\n43787 (2024).\n5. Ying, P., Li, B., Ma, M. Enhancing the hardness of diamond through twin\nrefinement and interlocked twins. Nature Synthesis, 4, 391-398 (2025).\n6. Huang Q., Yu D., Xu B., et al. Nanotwinned diamond with unprecedented hardness\nand stability. Nature, 510 (7504), 250-253 (2014).\nWe have added relevant references and following discussions in the revised\nmanuscript to better contextualize the relationship between NCD, UNCD, NTD, and\nour stacking-fault-dominated diamond structure. Please let us know if any additional\nreferences or comparisons would be helpful.\n“Nanocrystalline 13, 14 and ultrananocrystalline diamond (UNCD) 15, 16 films,\ncharacterized by nanoscale grain sizes, typically achieve enhanced hardness through\ngrain boundary strengthening mechanisms.” (Please see Page 3 in the revised\nmanuscript)\n13. Gruen, D. M. Nanocrystalline Diamond Films. Annual Review of Materials Science.\n29, 211-259 (1999).\n14. Sumiya H, I. T. Microstructure and Mechanical Properties of High- Hardness\nNano-Polycrystalline Diamonds. SEI Technical Review-English Edition. 66, 85 (2008).\n15. Lin, C. et al. Formation of ultrananocrystalline diamond films with nitrogen\naddition. Diam. Relat. Mat. 20, 380-384 (2011).\n16. Tu, J. et al. Microstructure control and mechanical properties of ultra-\nnanocrystalline diamond films. Ceram. Int. 50, 43780-43787 (2024). (Please see Page\n22 in the revised manuscript)\n3. The use of nanocrystalline diamond films in cutting tools is already well-established.\nA critical review of existing applications and comparative analysis should be added to\nenhance the relevance of the current results.\nResponse 3: We agree that nanocrystalline diamond (NCD) films have been widely\nexplored for cutting tool applications due to their excellent hardness, wear resistance,\nand chemical inertness. In the revised manuscript, we have added a brief review of\nrepresentative studies on NCD and ultrananocrystalline diamond (UNCD) films applied\nin machining and cutting tools. Compared with traditional NCD coatings, our stacking-\nfault-dominated diamond films not only exhibit comparable or superior hardness but\nalso demonstrate significantly improved thickness uniformity and scalability (up to 5\ninches), which are essential for industrial applications. We have also discussed the\npotential of our freestanding wafers to be integrated onto various tool surfaces, offering\nadvantages in both performance and processing flexibility. To enhance the contextual\nrelevance of the content, we have added relevant references and comparative analyses\nin the revised version as follows.\n“Owing to these mechanical advantages, both NCD and UNCD coatings have been\nwidely employed in cutting and machining tools for their high hardness, excellent wear\nresistance, and low friction performance 17-20. However, the mechanical performance of\nthese films is often constrained by factors such as limited thickness and grain-\nboundary-related degradation.” (Please see Page 3 in the revised manuscript)\n“17. Xia, Y. et al. Application of Nano-Crystalline Diamond in Tribology. Materials.\n16, 2710 (2023).\n18. Auciello, O. Science and technology of a transformational multifunctional\nultrananocrystalline diamond (UNCDTM) coating. Functional Diamond. 2, 1-24 (2022).\n19. Auciello, O. & Sumant, A. V. Status review of the science and technology of\nultrananocrystalline diamond (UNCD™) films and application to multifunctional\ndevices. Diam. Relat. Mat. 19, 699-718 (2010).\n20. Hu, J., Chou, Y. K., Thompson, R. G., Burgess, J. & Street, S. Characterizations of\nnano-crystalline diamond coating cutting tools. Surface and Coatings Technology. 202,\n1113-1117 (2007).” (Please see Page 22 in the revised manuscript)\n4. Please verify if “3 mm” refers to the actual thickness of the grown diamond. This is\nunusually high for CVD-grown films and requires confirmation or correction, while\nthinner samples were later tested.\nResponse 4: We have previously conducted extensive investigations and accumulated\nsignificant experience in the growth of thick diamond films (thickness > 2 mm) 1, 2. We\nconfirm that the 3 mm thickness refers to the actual diamond wafer obtained through a\ntwo-step fabrication process, which includes a conventional diamond layer and an\nultrahard diamond layer. This process is designed for the growth of large-area, high-\nthickness ultrahard diamond. In the first step, we employed a direct current arc plasma\njet chemical vapor deposition (DC Jet CVD) method to grow a 5-inch diameter, 3.5\nmm-thick free-standing polycrystalline diamond wafer on a graphite substrate. This\nmethod was chosen for its relatively high growth rate, which is suitable for large-scale\nbase wafer production. The as-grown wafer was then mechanically grinded and\npolished to a thickness of ~2.8 mm with a smooth surface. In the second step, this ~2.8\nmm-thick diamond wafer was used as a substrate in a 915 MHz microwave plasma\nCVD (MP-CVD) system. By applying a periodic nitrogen-doping strategy, an ultrahard\ndiamond layer with a thickness of approximately 200 μm was deposited on top. The\nhardness test is carried out on the surface of the ultra-hard diamond layer after polishing.\nTo avoid misunderstanding regarding the thickness and growth methodology, we have\nadded a clarification of this two-step process in the revised manuscript as follows:\n“A two-step growth strategy was employed to fabricate large-size, high-thickness free-\nstanding diamond wafers. In the first step, a 5-inch-diameter, 3.5 mm-thick free-\nstanding polycrystalline diamond wafer was deposited on a graphite substrate using a\ndirect current arc plasma jet chemical vapor deposition technique, which offers a high\ngrowth rate suitable for thick diamond synthesis. The wafer was then mechanically\ngrinded and polished to a final thickness of ~2.8 mm with a smooth surface. In the\nsecond step, an ultrahard diamond layer (~200 μm thick) was deposited on the polished\nwafer using a 915 MHz MP-CVD system under periodically modulated nitrogen doping,\nunder the following growth conditions: microwave power of 11-12 kW, chamber\npressure of 11-11.5 kPa, and CH₄/H₂/N₂/O₂ flow rates of approximately\n45/900/15/1.5 sccm. The temperature change caused by the gas cycle change is shown\nin Fig.1d.” (Please see Page 18 in the revised manuscript)\n1. Liu P, et al. The uniform and robust 8-inch CVD diamond plate generated by dual-\nmagnetic field controlled DC Jet CVD. Diamond and Related Materials, 141,\n110670 (2024).\n2. Huang Y, et al. The 7-in. freestanding diamond thermal conductive film fabricated\nby DC arc Plasma Jet CVD with multi-stage magnetic fields. Diamond and Related\nMaterials, 122, 108812 (2022).\n5. The statement “can be additionally grown on surfaces of other complicated\nengineering tools” is unclear. In microwave plasma systems, the plasma is typically\nlocalized in a spherical or disc shape, which limits uniform coating on complex 3D\ngeometries, especially along the Z-axis. Please clarify or revise this statement.\nResponse 5: The intention of our original statement was to highlight that the cyclic\nnitrogen-doping CVD process we developed is applicable not only to flat substrates,\nbut also to commonly used tool surfaces, which prompted us to explore its use in\ncoating cutting tools. Due to the localized nature of the plasma region in conventional\nmicrowave plasma systems, diamond growth is typically more suitable for planar\nsubstrates. For complex 3D geometries, plasma concentration and non-uniform\ndistribution can lead to uneven diamond coatings. However, for tools with relatively\nsimple shape and commonly used tools (Fig. 2e) uniform deposition can still be\nachieved by carefully designing the molybdenum holder to optimize substrate\nplacement and plasma exposure. The original manuscript may not have conveyed this\npoint clearly. Therefore we have revised relevant sentences and included the following\nclarification in the revised manuscript.\n“This periodic nitrogen-doping process can not only be applied to large-scale planar\nsubstrates, but can also be extended to commonly used tools, where uniform diamond\ncoatings can be achieved by optimizing the holder design and plasma exposure\nconditions.” (Please see Page 9 in the revised manuscript)\n6. The manuscript presents nitrogen as the sole effective parameter for stacked-diamond\nformation. Please comment on the potential role of other gases, such as argon-based\nplasmas or methane-oversaturated gas mixtures with/without nitrogen, which are also\nreported to affect the growth mechanisms in a similar manner.\nResponse 6: In our study, nitrogen was selected as the key doping species due to its\nwell-established role in modifying diamond growth kinetics and facilitating stacking\nfault formation by altering the local bonding environment and surface energy. However,\nwe fully agree that other gas species, such as argon or methane-rich gas mixtures,\ncan also significantly influence diamond growth mechanisms. Previous reports have shown\nthat argon-based plasmas can enhance secondary nucleation and promote\nnanocrystalline or ultra-nanocrystalline diamond growth through increased ion\nbombardment and reduced atomic hydrogen density, which may indirectly lead to\ndefect formation. Similarly, CH₄-oversaturated environments, especially in\ncombination with nitrogen, can enhance non-equilibrium growth conditions that favor\ndefect generation and nano-structuring. We have added the relevant discussions in the\nrevised manuscript as follows.\n“In the context of polycrystalline diamond film fabrication, previous studies have shown\nthat incorporating argon or employing methane-rich gas mixtures can significantly\ninfluence diamond growth kinetics and defect formation 27-30.” (Please see Page 5 in\nthe revised manuscript)\n“27. Martyanov, A. et al. Synthesis of Polycrystalline Diamond Films in Microwave\nPlasma at Ultrahigh Concentrations of Methane. Coatings. 13, 751 (2023).\n28. Chowdhury, S., Hillman, D. A., Catledge, S. A., Konovalov, V. V. & Vohra, Y. K.\nSynthesis of ultrasmooth nanostructured diamond films by microwave plasma chemical\nvapor deposition using a He/H /CH /N gas mixture. J. Mater. Res. 21, 2675-2682\n2 4 2\n(2006).\n29. Jiang, J. H. & Tzeng, Y. H. Mechanisms of suppressing secondary nucleation for\nlow- power and low-temperature microwave plasma self-bias- enhanced growth of\ndiamond films in argon diluted methane. AIP Adv. 1, 42117 (2011).\n30. Gruen, D. M., Liu, S., Krauss, A. R., Luo, J. & Pan, X. Fullerenes as precursors for\ndiamond film growth without hydrogen or oxygen additions. Appl. Phys. Lett. 64, 1502-\n1504 (1994).” (Please see Page 23 in the revised manuscript)\nIn addition, we have also elaborated on the role of each gas used in this work and\nthe dynamic regulation mechanism of the plasma environment in the revised Discussion\nsection. Specifically, H₂ serves as the main carrier gas, ensuring stable plasma discharge\nand etching non-diamond phases via H radicals to retain sp³-bonded diamond. CH₄\nprovides the carbon source for diamond growth. N₂ acts as a doping gas to disrupt the\natomic stacking sequence and induce stacking faults. A small amount of O₂ helps\noxidize non-diamond carbon, improving crystalline quality. The pulsed introduction of\nN₂ creates periodic perturbations in the plasma composition and substrate temperature,\nresulting in localized non-equilibrium growth conditions. These fluctuations promote\nthe formation of stacking faults and other defects, thereby enhancing the mechanical\nproperties of the diamond films. We have added relevant content in the Methods and\nDiscussion sections in the revised manuscript as follows.\n“During diamond deposition, H₂ maintains a stable plasma and etches the non-\ndiamond phase, while CH₄ provides a carbon source. N₂ doping disrupts the atomic\narrangement, while a small amount of O₂ improves the crystal quality by oxidizing the\nnon-diamond carbon.” (Please see Page 18 in the revised manuscript)\n“The periodic introduction of nitrogen perturbs the plasma and temperature\nenvironment, inducing localized non-equilibrium conditions that promote stacking fault\nformation.” (Please see Page 13 in the revised manuscript)\n7. Provide the growth rate corresponding to the reported 3 mm thick diamond layer.\nThis would substantiate claims regarding the high growth efficiency.\nResponse 7: We thank you for pointing out the importance of clarifying the growth rate\nto support the high deposition efficiency. In our two-step process for fabricating large-\narea, high-thickness ultrahard diamond wafers, the initial growth was performed using\nthe DC Jet CVD method, which provides a significantly higher deposition rate than\nconventional MP-CVD systems. The diamond growth rate ranged from 10 to 15 μm/h\n1 from the center to the edge of the wafer, enabling efficient deposition of thick diamond\nfilms. Subsequently, an ultrahard diamond layer was deposited using MP-CVD with\nprecisely controlled periodic nitrogen doping. In this process, the growth rate was\napproximately 410 nm/h, corresponding to 22.1 nm per gas-switching cycle. This\nslower, finely controlled deposition ensures the formation of a high-density stacking\nfault structure, which is critical for achieving ultrahigh hardness.\nThis two-step strategy combines the advantages of rapid thick film growth and\nprecise microstructural engineering to achieve superior mechanical performance,\nthereby meeting the demands of diverse application scenarios. We have added the\ncorresponding growth rates to the revised manuscript.\n1. Liu P, et al. The uniform and robust 8-inch CVD diamond plate generated by dual-\nmagnetic field controlled DC Jet CVD. Diamond and Related Materials, 141,\n110670 (2024).\n8. A more detailed discussion of the XRD peak intensities and the corresponding\ncrystallographic information is needed.\nSimilar detailed analysis should be applied to other techniques used (e.g., Raman\nspectroscopy). Spectral features and peak assignments should be scientifically\ninterpreted rather than descriptively reported.\nResponse 8: To provide a more rigorous interpretation of the spectral features, we have\nrevised and expanded the discussion of the XRD and Raman spectroscopy results in the\nrevised manuscript as follows.\n“The XRD pattern of the PCD film (Fig.1f) exhibits a strongest (111) diffraction peak\nat 44° (predominantly exposed crystal plane), along with weaker (220) and (311) peaks\nand a very faint (400) peak. The narrow full width at half maximum of the (111) peak\nsuggests high crystallinity.” (Please see Page 5-6 in the revised manuscript)\n“The Raman spectrum shows three characteristic peaks at 1195 cm⁻¹, 1332.5 cm⁻¹,\nand 1520 cm⁻¹. The sharp peak at 1332.5 cm⁻¹ corresponds to the first-order Raman\nmode of sp³-bonded diamond. The 1195 cm⁻¹ peak is associated with defect-related or ultra-\nnanocrystalline diamond features. The 1520 cm⁻¹ peak may be attributed to nitrogen-\ninduced local lattice distortions, distinct from the typical G band of sp² carbon.”\n(Please see Page 6 in the revised manuscript)\n9. The comparison of hardness should be presented with better clarity: a) Specify the\nsubstrate sizes, especially for SCD samples and b) since hardness tests are performed\nat the micro-scale, additional commentary on macro-scale hardness implications should\nbe included and finally, c) it is not clear the presented hardness values of PCD and SCD\nsubstrates were measured by the authors using the same testing methods as those used\nfor the current samples (please, clarify or measure PCD/SCD in your setup), while\nreliance on external literature values may lead to inconsistent or misleading\ncomparisons.\nResponse 9: We thank you for the insightful comment regarding the hardness\ncomparison.\na) The substrate sizes used in our study have now been specified in the revised\nmanuscript. For clarity, the polycrystalline diamond (PCD) substrate measured 5 inches\nin diameter and 2.8 mm in thickness, while the single-crystal diamond (SCD) sample\nwas a 5×5×1 mm3 colorless CVD-grown diamond plate.\nb) Thank you for your insightful comment and valuable suggestion. In our study, the\nVickers hardness tests were conducted under load of 1 kgf (9.8 N) to obtain asymptotic\nhardness, producing indentation diagonals of approximately 9-10 μm, which fall within\nthe mesoscopic scale. While Vickers hardness measurements at this scale primarily\nreflect local mechanical properties, the mesoscopic scale serves as a bridge between the\nmicro- and macro-scale behavior, and is highly relevant to the material’s bulk\nperformance-especially in wear-resistant applications where ultrahard diamonds are\ntypically used. The ultrahard diamond we prepared has excellent wear resistance (see\nFig. 2h and Supplementary Fig. 5).\nNotably, the HSD wafers in our study exhibit a uniform microstructure and\nultrahigh hardness across large areas (up to 5 inches in diameter), indicating excellent\nscalability of the mechanical properties. Moreover, the high-density stacking fault\nnetwork formed via the cyclic nitrogen doping process is distributed throughout the\nentire wafer, which helps maintain consistent hardness and structural integrity both\nlocally and over larger scales.\nTo ensure the reliability of the measurements, hardness testing was performed at\nmultiple randomly selected positions across each sample. The consistent results\nobtained confirm the homogeneity of the material and support the relevance of meso-\nscale Vickers hardness as a meaningful indicator of macro-scale performance. To clarify\nthe connection between the micro-/meso-scale hardness values and their macroscopic\nimplications, we have added the following discussion in the revised manuscript.\n“Vickers hardness tests were conducted under 1 kgf (9.8 N) load, producing\nindentations with diagonal lengths of approximately 9-10 μm, corresponding to the\nmesoscopic scale. Multiple indentations were randomly distributed across the sample\nsurface to ensure measurement accuracy and uniformity.” (Please see Page 8 in the\nrevised manuscript)\n“This meso-scale hardness measurement reliably reflects the macroscopic mechanical\nperformance of the material, especially given the relatively uniform microstructure\n(Supplementary Fig. 3) and little variation in hardness observed across the wafer.”\n(Please see Page 8 in the revised manuscript)\nc) The hardness values of both the PCD and SCD substrates were measured using the\nsame indentation equipment and test protocols as for the HSD wafers (see Fig. R2),\nensuring consistency across all datasets. We have clarified this point in the revised text\nto\navoid any confusion, and we avoided relying on literature values to minimize\ninconsistencies or misinterpretation.\nFig. R2 a Morphology of PCD indentation. b Morphology of SCD indentation. a , a\n1 2\nare the lengths of the two diagonals of the indentation. b , b are the lengths of the two\n1 2\ndiagonals of the cracks.\nWe hope this clarification adequately addresses your concerns, and we appreciate\nthe valuable suggestion to improve the accuracy and transparency of our comparisons.\n10. The growth of diamond films on 5-inch planar substrates using microwave plasma\nCVD at 915 MHz is not currently a technological breakthrough. Phrases such as “5 inch\nimpressive” are considered overstated and should be omitted or rephrased.\nResponse 10: We agree that the description of the 5-inch substrate growth should be\npresented with more scientific objectivity. In the revised manuscript, we have removed\nsubjective terms such as “impressive” and rephrased the relevant statements to provide\na more accurate and balanced representation of the technological context.\n11. The information regarding direct HSD growth on WC inserts is potentially\nmisleading due to the catalytic graphitization effect of cobalt in WC. Either elaborate\non any barrier layers or initial pretreatments of WC used, or revise this statement\naccordingly.\nResponse 11: The statement regarding “direct deposition on various material surface”\nmay have been misleading and has now been revised. Our intention was to highlight\nthe applicability of the cyclic nitrogen doping strategy to various suitable substrates.\nIndeed, the presence of cobalt in WC can catalyze the graphitization of diamond and\nhinder stable growth. In our experiments, the WC tools were first subjected to surface\npretreatment to remove cobalt, and then a thin nanocrystalline diamond interlayer was\ndeposited to serve as a transition layer and improve nucleation compatibility. We\ndeleted the word \"direct\" and added the relevant content in the revised manuscript as\nfollows:\n“After cobalt removal and deposition of a nanocrystalline diamond transition layer, the\nHSD coating was successfully grown on the surface of the tungsten carbide end mill.”\n(Please see Page 9 in the revised manuscript)\n12. Toughness values calculated using elastic modulus may be overestimated.\nNanocrystalline diamond materials can exhibit values 10-25% lower than theoretical\nestimations. Please revise and discuss the limitations of this method.\nResponse 12: Due to the presence of grain boundaries, micropores, and\ncrystallographic anisotropy, the elastic modulus of polycrystalline diamond films is\ngenerally lower than that of single-crystal diamond, typically in the range of 850-1050\nGPa, which is about 10-20% lower than that of the ideal single crystal diamond.\nTherefore, the fracture toughness values calculated by theoretical formulas using an\nelastic modulus of 1000 GPa may be overestimated.\nIn our study, to obtain a reliable value of the elastic modulus for the ultrahard\ndiamond film, we performed surface acoustic wave (SAW) dispe𝐸𝐸rsion analysis by\nmeasuring the phase velocity as a function of frequency, as shown in Fig. R3. Using\nRayleigh wave approximations, the elastic modulus of the HSD-1 sample was\ndetermined to be approximately 1033 GPa. This result confirms that the toughness\ncalculated using an assumed modulus of 1000 GPa in the manuscript is not\noverestimated. It is worth noting that commonly used values for diamond toughness\ncalculations are 1000 GPa 1 or 1100 GPa 2. To ensure consistency and facilitate\ncomparison with previously reported results on diamond materials, we adopted 1000\nGPa as the elastic modulus in our calculations.\nFig. R3 Phase velocity of the surface acoustic wave versus the frequency for HSD-1\nsample\n1. Huang Q, Yu D, Xu B, et al. Nanotwinned diamond with unprecedented hardness\nand stability. Nature, 510(7504), 250-253 (2014).\n2. Yue Y, Gao Y, Hu W, et al. Hierarchically structured diamond composite with\nexceptional toughness. Nature, 582(7812), 370-374 (2020).\n13. Some labels in Figure 2 are not clearly linked to specific experimental results or\nreferences. Please revise for better clarity and traceability.\nResponse 13: We have revised Figure 2d to improve clarity and traceability.\nSpecifically, all labels have been checked and updated to ensure they are clearly\nassociated with corresponding experimental results or literature references. The\nfollowing table R2 provides the specific hardness and sample size information\ncorresponding to each reference. These modifications enhance the readability and\nscientific rigor of the figure, as shown in the revised version of the manuscript (see\nFigure 2 and its updated caption).\nTable R2 Comparison of hardness and sample size of different types of diamonds\nSample\nHardness/GPa Reference\nsize/mm\nSCD 79.7 0.8-1 Suprun O.M., Crystals, 7, 369 (2017).\nYan C. S., Physica Status Solidi (a) 201,\n160 4.5\n4, R25-R27 (2004).\nLiang Q., Journal of Physics: Condensed\n78±12, 125 5\nMatter, 21, 364215 (2009).\n115.4, 126.4, 103.4, 111.7, 117.1, 96.7, Sumiya H., Diamond and Related\n5-6\n98.3, 101.9, 83.0, 92.4, 84.3, 94.8, 92.9 Materials, 1997,6 (12) 1841-1846.\n110-140 1 Irifune, Nature 421, 599-600 (2003).\nLu J., Diamond and Related Materials,\n125 3\n77:41-45 (2017).\nLian M., International Journal of\nPCD 124.5, 121.3,120.7, 120.5, 114.6, 113.3 4.2 Refractory Metals and Hard Materials\n118, 106490 (2024).\nLi Q., Scientific Reports, 10: 22020\n120, 11\n(2020).\nLiu J., Applied Physics Letters, 112 (6):\n120.5±4.2 11\n061901 (2018).\nSumiya H., Diamond & Related\n105, 95, 138, 141, 145, 131 1\nMaterials, 13, 1771-1776 (2004).\n128, 138, 120, 140, 113, 139, 97, 119, 70, Sumiya H., Journal of materials research,\n1-2\n74, 95, 112, 66, 86, 70, 85, 93, 104 2007, 22(8): 2345-2351.\nNPD\nZhao B., Nanoscale Advances, 2: 691\n90 3\n(2020).\nYao M. G., Advanced Functional\n105, 145, 130, 60, 65, 84, 118, 83 1.5\nMaterials, 33, 2214696 (2023).\n200 2 Huang Q., Nature, 2014, 510, 250-253.\nNt-\ndiamond Ying P., Nature Synthesis, 4, 391-398\n276 1\n(2025).\n“Fig. 2 d Statistical plot of diamond hardness as a function of sample diameter (The\nsquare shape of the samples according to its longest side as the diameter).” (Please\nsee Page 11 in the revised manuscript)\n14. The concept of cyclic gas-phase doping has been previously explored, particularly\nfor phosphorus doping. Relevant studies should be cited, and the novelty of the current\ndoping method should be more clearly positioned.\nResponse 14: Indeed, cyclic or pulsed doping methods have been previously employed,\nnotably in phosphorus doping processes for semiconductors and CVD-grown diamond.\nFor instance, cyclic or pulsed gas-phase doping techniques have been reported to\nenhance doping efficiency and modulate impurity incorporation in materials such as\nsilicon and diamond. These techniques typically involve low-frequency modulation of\nthe dopant gas flow to tune the electronic or optical properties. We have now cited\nrepresentative studies in the revised manuscript to better acknowledge this prior work.\nIn contrast, the novelty of our method lies in applying a high-frequency, short-period\nnitrogen doping cycle during MP-CVD growth to engineer stacking fault structures in\ndiamond. Unlike conventional electronic-grade doping, our approach does not aim to\nincorporate nitrogen as a shallow donor but instead utilizes transient nitrogen\nincorporation to locally destabilize the lattice and promote the formation of non-\nequilibrium planar defects. This strategy enables the retention of a dense stacking-fault\nnetwork that enhances mechanical hardness—fundamentally distinct in purpose and\noutcome from previous doping approaches. To better clarify this distinction, we have\nadded relevant references and explicitly discussed the novelty and mechanism of our\nnitrogen doping strategy in the revised manuscript as follows.\n“Although cyclic gas-phase doping has been previously applied to modulate electrical\nproperties of diamond, such as in phosphorus doping strategies34-36, our work uniquely\napplies high-frequency, short-duration nitrogen pulses to induce defects during\ndiamond growth, introducing a new pathway for defect-driven mechanical\nenhancement.” (Please see Page 5 in the revised manuscript)\n“34. Lobaev, M. A., Radishev, D. B., Gorbachev, A. M., Vikharev, A. L. & Drozdov,\nM. N. Investigation of microwave plasma during diamond doping by phosphorus using\noptical emission spectroscopy. Physica Status Solidi (a). 216, 1900234 (2019).\n35. Kociniewski, T. et al. n‐type CVD diamond doped with phosphorus using the\nMOCVD technology for dopant incorporation. Physica Status Solidi (a). 203, 3136-\n3141 (2006).\n36. Lloret, F. et al. High phosphorous incorporation in (100)-oriented MP CVD\ndiamond growth. Diamond Related Materials. 133, 109746 (2023).” (Please see Page\n23 in the revised manuscript)\nLastly, we thank you again for your insightful review and suggestions that improved\nour manuscript.\nReviewer #3 (Remarks to the Author):\nThis manuscript presents a novel approach for synthesizing large-scale ultrahard\ndiamond wafers using a customized microwave plasma chemical vapor deposition\n(MP-CVD) method with high-frequency pulsed local non-equilibrium growth. The\nauthors report achieving diamond wafers up to 5 inches in diameter with exceptional\nhardness (~208.3 GPa) through periodic nitrogen doping, which they attribute to the\nformation of a three-dimensional interlocked stacking fault network. This manuscript\npresents potentially significant advances in ultrahard diamond synthesis with\nimpressive scale and claimed properties. However, several critical aspects require\nfurther clarification and validation before publication. I recommend major revision to\naddress the concerns outlined in the following, particularly regarding the hardness\nmeasurement methodology, stacking fault characterization, and stacking fault energy\ncalculations. If these concerns can be adequately addressed, this work would represent\na valuable contribution to the field of ultrahard materials and diamond synthesis.\nResponse: We first of all thank you very much for your positive assessment of our work\nand for recognizing the potential significance of our approach to synthesizing large-\nscale ultrahard diamond wafers. We fully acknowledge the importance of the concerns\nyour raised, particularly with respect to the hardness measurement methodology,\nstacking fault characterization, and stacking fault energy calculations. In the revised\nmanuscript, we have made substantial efforts including some key improvements to\naddress these issues. We believe that with your help, these revisions have substantially\nstrengthened the manuscript and improved its clarity, rigor, and scientific impact.\nDetailed responses to each specific comment are provided below.\n1. a) The reliability of the hardness measurements is questionable based on the evidence\npresented. The Vickers indentation shown in Fig. 2a inset displays obvious features\nwith ridges extending beyond the marked area, suggesting potential overestimated\nhardness. b) The reported PCD hardness (~100 GPa) is significantly higher than typical\nvalues (~70 GPa), raising questions about the calibration of the measurement system.\nc) Additionally, the authors don't specify which crystallographic plane of the single-\ncrystal diamond was measured, which is critical since diamond hardness varies\nsignificantly with crystallographic orientation. The reported fracture toughness for SCD\n(~9.1 MPa·m^(1/2)) is also unusually high compared to typical values (less than 5\nMPa·m^(1/2)), further suggesting potential measurement inconsistencies.\nResponse 1: Thank you for your constructive comments regarding the hardness\nmeasurements and related characterization. a) The inset in Fig. 2a was intended to\nillustrate the indentation morphology, and the white dashed diamond outline was a\nschematic indicator rather than a measurement reference. In the actual hardness\ncalculations, we determined the Vickers hardness and fracture toughness based on\nprecise measurements of the full diagonal lengths of the indentation (a₁ and a₂) and the\nradial crack lengths (b₁ and b₂) extending from the corners, as illustrated in the updated\nimage now included in Fig. 2a. To avoid any potential misunderstanding, we have\nreplaced the original inset with a revised figure that clearly marks the measured features\nused in the calculations.\n“Fig.2 The inset shows the typical indentation morphology of HSD-1 and the diagonal\nlengths of the indentation and crack.” (Please see Page 11 in the revised manuscript)\nAdditionally, we would like to clarify a common source of deviation in Vickers\nhardness measurements. The primary source of error arises from determining the\naccurate diagonal lengths of the indentation. The ideal diagonal length is the distance\nbetween the two vertices of the indentation on the reference plane of the sample surface\n1, 2 (d in Fig. R4). However, in practice, small raised ridges can form around the\n0\nindentation due to plastic deformation, leading to a slightly larger measured diagonal\n(d) than the actual value. This typically results in a slight underestimation of the true\nhardness, not overestimation. Therefore, the hardness values reported in this work are\nnot overestimated and may in fact be slightly lower than the intrinsic values.\nFig. R4 a Indentation morphology map of the HSD-1 sample formed under a load of\n9.8 N; b the profile lines a and a .\n1 2\n1. Xu B., Tian Y. J. Ultrahardness: Measurement and Enhancement. The Journal of\nPhysical Chemistry C, 119 (10), 5633-5638 (2015).\n2. Xu B., Tian Y. J. High pressure synthesis of nanotwinned ultrahard materials.\nActa Physica Sinica, 66 (3): 036201 (2017).\nb) About the hardness of PCD, the hardness of diamond can vary widely (typically\nbetween 60-120 GPa) 3, 4, depending on factors such as crystallographic orientation and\ngrain size. The PCD substrates used in this work were synthesized using the DC Jet\nCVD method, which promotes columnar crystal growth. The sidewalls of these\ncolumnar grains contain a high density of small grains and twin boundaries, both of\nwhich are known to contribute to hardness enhancement 5. To ensure the reliability of\nour results, we performed multiple repeated hardness measurements across different\nareas of the PCD sample. The measured values remained consistent with those reported\nin the manuscript, indicating that the hardness testing procedure was correctly\nconducted and the reported results are valid. The hardness is calculated based on the\ndiagonal length of the indentation. The following Fig. R5 shows the typical indentation\nmorphology of the tested PCD substrate.\nFig. R5 Morphology of PCD indentation. a , a are the lengths of the two diagonals of\n1 2\nthe indentation. b , b are the lengths of the two diagonals of the cracks.\n1 2\n3. Nie, A., Zhao, Z., Xu, B. et al. Microstructure engineering in diamond-based\nmaterials. Nat. Mater. (2025).\n4. Brazhkin, V., Dubrovinskaia, N., Nicol, M. et al. What does 'harder than diamond'\nmean? Nature Mater 3, 576-577 (2004).\n5. An K., Chen L., Yan X., et al. Fracture behavior of diamond films deposited by DC\narc plasma jet CVD. Ceramics international, 44 (11): 13402-13408 (2018).\nc) Concerning the crystallographic orientation, we thank you for pointing this out.\nFor single-crystal diamond (SCD), the indentation tests were performed on the {100}\nsurface. We have included this critical detail in the revised manuscript to ensure clarity\nand reproducibility.\nIn addition, we appreciate your comment regarding the reported fracture toughness\nof SCD. As documented in the literature, the fracture toughness of SCD varies\nsignificantly depending on its type and impurity content. For instance, natural type\nIa/IIa diamonds and colorless CVD SCD typically exhibit fracture toughness around\n8 ± 4 MPa·m1/2. In contrast, type Ib yellow diamonds synthesized without N₂ show\nhigher toughness values of approximately 10 ± 2 MPa·m1/2 6, 7, while nitrogen-doped\nSCD can reach 13.8 ± 2.7 MPa·m1/2 8. Polycrystalline diamond has poor toughness due\nto the influence of grain boundaries, which is generally 5±2 MPa·m1/2 9-11.\nThe SCD used in our study is a colorless CVD-grown diamond, and the measured\ntoughness of 9.1 MPa·m1/2 falls within the expected and reasonable range for such\nmaterial. The fracture toughness was calculated based on the crack lengths observed in\nthe Vickers indentations. For clarity, the representative images of the indentation used\nfor this calculation are shown in Fig. R6.\nFig. R6 Morphology of single-crystal diamond indentation. a , a are the lengths of\n1 2\nthe two diagonals of the indentation. b , b are the lengths of the two diagonals of the\n1 2\ncracks.\n6. Liang Q., Yan C., Meng Y., et al. Enhancing the mechanical properties of single-\ncrystal CVD diamond. Journal of Physics: Condensed Matter, 21 (36), 364215\n(2009).\n7. Yan C., Mao H., Li W., et al. Ultrahard diamond single crystals from chemical\nvapor deposition. physica status solidi (a), 201 (4), R25-R27 (2004).\n8. Zhao Y., Tu J., Chen L., et al. Toughness enhancement of single-crystal diamond\nby the homoepitaxial growth of periodic nitrogen-doped nano-multilayers.\nInternational Journal of Minerals, Metallurgy and Materials, 30 (4), 766-771\n(2023).\n9. Drory M., Gardinier C., Speck J. Fracture toughness of chemically vapor‐\ndeposited diamond. Journal of the American Ceramic Society, 74 (12), 3148-3150\n(1991).\n10. Drory M., Dauskardt R., Kant A., et al. Fracture of synthetic diamond. Journal of\nApplied Physics, 78 (5), 3083-3088 (1995).\n11. McNamara D., Carolan D., Alveen P., et al. Effect of loading rate on the fracture\ntoughness and failure mechanisms of polycrystalline diamond. International\nJournal of Refractory Metals and Hard Materials, 60, 1-10 (2016).\n2. a) The stacking fault characterization presented in the manuscript requires substantial\nimprovement. Many of the TEM images lack sufficient resolution and clarity to\nunambiguously identify stacking fault structures. The quasi-periodic structures visible\nin Fig. 3b warrant higher-resolution atomic-scale characterization to confirm their\nnature. b) While the authors report a stacking fault density (usually defined as SF\narea/volume) of 4.3×10^12 cm^-2, the statistical sampling methodology is not clearly\ndescribed, raising questions about the representativeness of this value. Given that the\nstacking fault network is presented as the primary mechanism for the enhanced\nhardness, more comprehensive and higher quality microstructural characterization is\nessential.\nResponse 2: a) We also believe that in-depth and detailed characterization and\ninterpretation of the stacking fault structures are critical to our work. Following your\nadvice, we have improved the quality of the HRTEM images, as shown below Fig. R7\n(Fig. 3b in the revised manuscript). The spatial resolution of the Cs-corrected HRTEM\nimage now reaches ~75 pm. Distinct streaks can be clearly observed in the FFT pattern\n1, which are indicative of stacking fault structures in the equivalent {111} planes.\nFig. R7 a HRTEM image of the HSD sample. b Corresponding FFT pattern showing\nthe spatial frequency of a.\nTo further verify the atomic structure of the stacking faults, Cs-corrected HAADF-\nSTEM imaging was conducted, which further confirms that the defect structures within\nour sample are indeed stacking faults, as shown below (Fig. R8).\nFig. R8 Atomic-resolution HAADF-STEM image confirming the typical diamond\nstacking fault structure in HSD-1.\nFor the quasi-periodic structures observed in Fig. 3b, we also performed detailed\natomic-resolution HAADF-STEM experiments (see Fig. R9), and found that these\nstructures result from the common overlapping SFs within diamond 2, 3. A detailed\nexplanation is provided below (also see Supplementary Fig. 13).\nFig. R9 Further analysis of the quasi-periodic structures inside the sample. a\nBright-field TEM image showing representative quasi-periodic structure. b Cs-\ncorrected HAADF-STEM enlarged view of the quasi-periodic structure (yellow box in\na), with the corresponding atomic model showing the local atomic arrangement (blue\nbox).\nSupplementary Fig. 13 Detailed explanation of quasi-periodic structure observed in\nthe HSD sample. a Schematic representation of TEM imaging for the FIB lamella of\nHSD sample (illustrated with conventional TEM mode as an example). It is evident that\nif the planar defect region in the HSD sample is very small (i.e., thickness along the\nelectron beam direction less than the total thickness of the FIB lamella), the TEM image\nwill display an overlap of two different structures: a perfect crystalline structure and a\nplanar defect structure. b Atomic-resolution HAADF-STEM images showcasing two\nexamples of overlap structure inside HSD samples, with the corresponding atomic\nmodels.\nb) Thank you very much for your valuable comments regarding the representativeness\nof the calculated stacking fault density value. As you pointed out, SF density can be\ndefined as SF area/volume, or it can be defined as SF number/area 4, 5.\nIn this work, we primarily adopted the latter definition (SF number/area) to\ncalculate SF density based on TEM images, as shown in Fig. R10, which involves\ncounting the number of stacking faults over a large area in bright-field TEM images.\nFor accurate counting, the contrast of the TEM images was inverted, and the intensity\nthreshold was adjusted to highlight the stacking fault regions (i.e., the white stripes in\nFig. R10b). Each white stripe was counted as one SF, allowing us to determine both the\ntotal number and the spacing of SFs. The SF density value of 4.3×1012/cm2 was\nobtained by averaging the measurement results from 10 TEM images taken from\ndifferent regions of the sample.\nFig. R10 Procedure for stacking fault density calculation. a Bright-field TEM image\nof HSD-1. b Contrast-adjusted image of a, in which the SFs are more clearly displayed,\nfacilitating accurate counting of the number of SFs.\nIn addition, we also followed your suggestion and recalculated the SF density\nusing the “SF area/volume” method. Here, we selected an HRTEM image along the\n<111> zone axis and directly measured the area of SFs on the {111} planes, as shown\nin Fig. R11. As a result, the SF density was determined to be 8.7 × 106/cm, which is\nvery close to the value obtained above (as the average length of the SFs is 24.7 nm (see\nSupplementary Fig. 8), corresponding to a SF density of 1.03 × 107/cm). It should be\nalso noted that this method may underestimate the actual SF density due to the possible\nSF overlaps. This consistency further confirms the reliability and accuracy of our SF\ndensity calculations.\nFig. R11 Procedure for SF density calculation using “SF area/volume” method. a\nHRTEM image of HSD-1 with the corresponding FFT pattern showing the <111> zone\naxis. b Contrast-adjusted image of a, with the SF region marked by white dashed lines\nfor measurement.\n1. Zhu, W, et al. Structural imperfections in CVD diamond films. Journal of Materials\nResearch, 4, 659-663 (1989).\n2. Tong K, et al. Structural transition and migration of incoherent twin boundary in\ndiamond. Nature, 626(7997), 79-85 (2024).\n3. Cook B, et al. Extensively microtwinned diamond with nanolaminates of\nlonsdaleite formed by flash laser heating of glassy carbon. Nano Letters, 23(22),\n10311-10316 (2023).\n4. Schreck M, et al. Charge carrier trapping by dislocations in single crystal diamond.\nJournal of Applied Physics, 127(12), 125102 (2020).\n5. Ichikawa K, et al. High crystalline quality heteroepitaxial diamond using grid-\npatterned nucleation and growth on Ir. Diamond and Related Materials, 94, 92-100\n(2019).\n3. There is a concerning ambiguity between the main text and the Methods section\nregarding which parameter is being varied to control nitrogen doping. The main text\nstates \"as the nitrogen doping duration per cycle decreased,\" but the Methods section\nindicates that the N₂ turn-on duration remained constant at 48 seconds. The variable\nparameter appears to be the methane gas flow duration after nitrogen introduction\n(parameter \"x\" in step 2), which should be clarified to ensure reproducibility of the\nprocess.\nResponse 3: The confusion is likely from our previous inaccurate explanation of the\nnitrogen doping process. In the periodic nitrogen doping process, each cycle consists of\nfive steps. In step 1, the nitrogen flow is turned on and maintained for 48 seconds to\nfully fill the microwave chamber and establish a nitrogen-rich equilibrium environment.\nImportantly, this step does not involve methane, and thus no diamond growth occurs\nduring this phase.\nActual nitrogen-doped diamond growth begins in step 2, when methane is\nintroduced alongside nitrogen. Therefore, the “nitrogen doping duration per cycle”\nmentioned in the main text specifically refers to the duration of step 2, during which\ndiamond growth occurs in the presence of nitrogen. We have thus revised the\nmanuscripts accordingly, to clarify this point and avoid such ambiguity.\n4. a) The stacking fault energy calculations presented in the manuscript require further\nvalidation. The calculated stacking fault energy values appear to be in the range of\nhundreds of mJ/m², which is much lower than typical values for diamond (~10 J/m²).\nComplete generalized stacking fault energy (GSFE) calculation results should be\nprovided as Supplementary Information to validate these values. b) Furthermore, the\nthermal stability of these high-density stacking fault structures is not addressed. Heat\ntreatment experiments would be valuable to understand how this microstructure might\nevolve under thermal loading, which is crucial for assessing the practical utility of these\nmaterials.\nResponse 4: a) We agree that the stacking fault energy (SFE) values are critical for\nunderstanding the defect behavior in diamond. In our manuscript, the calculated SFE\nvalues are indeed in the range of hundreds of mJ/m², which are consistent with previous\nfirst-principles studies. For instance, P. J. H. Denteneer 1 calculated the intrinsic and\nextrinsic stacking fault energies for diamond to be approximately 300 mJ/m2 and 253\nmJ/m2, respectively. These values validate the magnitude of our results. Moreover, we\nemphasize that the SFE values reported in our work were obtained through generalized\nstacking fault energy (GSFE) calculations, which describe the energy per unit area\nrequired to rigidly shift two parts of a crystal relative to each other along a given slip\nsystem. As defined by Yang et al.2, the GSFE is given by:\n,\n𝐸𝐸𝑑𝑑−𝐸𝐸0\nwhere, E and E denote the total ener𝛾𝛾g 𝑑𝑑 y =of th𝐴𝐴e0 system before and after the slip relative\n0 d\ndisplacement vector d, respectively, and A is the area of the stacking fault.\nvalues, confirming the correctness of our methodology. We further observed a reduction\nin GSFE after nitrogen doping, indicating that nitrogen facilitates stacking fault\nformation by reducing the energy barrier. This observation is consistent with the\nincreased stacking fault density seen in nitrogen-doped samples.\n1. Denteneer P J H. Stacking-fault energies in silicon, diamond, and silicon carbide.\nMRS Online Proceedings Library, 141(1), 343-348 (1988).\n2. Yang B, et al. Strengthening and toughening by partial slip in nanotwinned\ndiamond. Carbon, 150, 1-7 (2019).\nb) Regarding the thermal stability of high-density stacking fault structures, we\nevaluated the thermal stability of the ultra-hard polycrystalline diamond film in air\nusing a thermogravimetric analyzer (TGA55) at a heating rate of 10°C/min. The results\nare shown in the figure below. The thermogravimetry (TG) curve shows that the initial\noxidation temperature of diamond is 755°C, and the differential scanning calorimetry\n(DSC) curve shows that the initial oxidation temperature of diamond is 760 °C. The\nthermal oxidation temperature is comparable to that of natural diamond (770 ℃ from\nTG or 720 ℃ from DSC) and greater than that of nanocrystalline diamond (~680 °C).\nIn addition, to address your concerns regarding the microstructural evolution under\nthermal loading, we performed detailed in situ heating TEM experiments on the HSD\nsample at atomic resolution (from room temperature to 1100 °C). As shown in\nSupplementary Fig. 7, the stacking fault structures exhibit excellent thermal stability,\nremaining intact up to 900 °C. Notably, degradation of the HSD sample begins from the\nperfect crystalline regions (yellow dashed lines), rather than from the stacking fault\nregions (indicated by red arrows). The annihilation of defects is observed at 1000 °C in\nthe TEM. The in situ results clearly demonstrate that the thermal stability of HSD is not\ncompromised by the stacking fault structures.\nThis result indicates that the stacking fault structure has excellent thermal stability.\nIn the revised version, we have added the data and discussions as following:\n“The thermal stability of the HSD-1 sample was characterized by thermogravimetry\ncurves measured in air. At a heat ingrate of 10 ℃/min, the onset oxidation temperature\nwas ~760 ℃ from the thermogravimetry (TG) curve and the differential scanning\ncalorimetry (DSC) curve (Supplementary Fig. 6). The thermal oxidation temperature is\ncomparable to that of natural diamond (770 °C from TG or 720 °C from DSC)1 and\ngreater than that of nanocrystalline diamond (~680 °C)56. This result indicates that the\nstacking fault structure has excellent thermal stability.” (Please see Page 10 in the\nrevised manuscript)\n“Supplementary Fig. 6 Typical thermal stability of HSD-1 sample. TG a and DSC b\ncurves were measured in air at a heating rate of 10 ℃/min.” (Please see Page 9 in the\nrevised Supplementary Information)\n“1. Huang Q., Yu D., Xu B., Hu W., Ma Y., Wang Y., et al. Nanotwinned diamond with\nunprecedented hardness and stability. Nature, 510 (7504), 250-253 (2014).\n56. Solozhenko, V. L., Kurakevych, O. O. & Le Godec, Y. Creation of nanostuctures by\nextreme conditions: high-pressure synthesis of ultrahard nanocrystalline cubic boron\nnitride. Adv. Mater. 24, 1540–1544 (2012).” (Please see Page 22 and 24 in the revised\nmanuscript)\n“In situ heating TEM observations (Supplementary Fig. 7) reveal that the presence of\nstacking faults does not compromise the thermal stability of the HSD diamond, as\ndegradation initiates from the perfect crystalline regions rather than the faulted areas.”\n(Please see Page 10 in the revised manuscript)\n“Supplementary Fig. 7 In situ heating TEM investigation of the thermal stability of\nstacking faults in HSD. a-l HRTEM image sequence showing the evolution of stacking\nfaults in HSD at elevated temperatures (from room temperature up to 1100 °C). The\ntemperature was increased in steps of 100 °C, with a holding time of 10 min at each\nstep. Stacking fault regions are indicated by red arrows, and the degraded regions are\noutlined with yellow dashed lines. The onset temperature of degradation is slightly\nhigher than that observed in TG and DSC measurements on bulk samples, which may\nbe attributed to the vacuum environment within the TEM.” (Please see Page 10 in the\nrevised Supplementary Information)\n5. The manuscript would benefit from additional language editing to improve clarity\nand readability. There are several instances of grammatical errors and awkward\nphrasing throughout the text.\nResponse 5: Thank you again for your careful review. Now with your comments and\nsuggestions, along with other reviewers’, we have thoroughly revised the manuscript\nand improved the overall clarity, grammar, and readability. The revised manuscript has\nalso been reviewed by a native English speaker for linguistic accuracy and polishing.\nWe believe the revised version communicates our findings more precisely and\neffectively. With that, we look forward to the publishing of this work can stimulate more\ninterests and efforts in developing advanced ultrahard diamond materials and\napplications."}, {"round": 2, "reviewer_comments": {"reviewer_1": "I am satisfied with the revision, recommend for publication now.", "reviewer_2": "The authors have thoroughly addressed all reviewer concerns. Taking into account their responses to other referees as well, the scientific integrity of the manuscript has been significantly improved, and the manuscript is recommended for publication in Nature Communications.", "reviewer_3": "The authors have made substantial efforts to address my initial concerns, and I find their work on synthesizing 5-inch ultrahard diamond wafers to be a notable achievement that merits publication with minor adjustments.\nHowever, I suspect the hardness values may still be overestimated. Based on the cited literature, AFM measurements typically overestimate hardness by about 20% compared to optical microscopy methods. Even with this correction, yielding approximately 180 GPa, this remains impressive for a 5-inch sample and should be acknowledged.\nRegarding the thermal stability discussion, there is a structural issue with the manuscript organization. The authors discuss stacking fault evolution during thermal treatment before they have properly identified and characterized the stacking faults through TEM analysis. This discussion should be repositioned to appear after the TEM characterization section where stacking faults are definitively identified and analyzed.\nMost concerning are the reported GSFE values, which appear much lower than recent literature including both their cited Carbon paper and a 2023 Physical Review B publication (10.1103/PhysRevB.108.014104). This discrepancy requires better explanation and comparison with current calculations.\nDespite these concerns, the work represents a valuable contribution to ultrahard materials research. The synthesis of large-scale diamond wafers with engineered stacking fault networks is a significant technological advancement. With these adjustments addressing measurement limitations, manuscript organization, and literature consistency, the manuscript should be acceptable for publication."}, "author_response": "Responses to Reviewers’ Comments\nInch-Scale Ultrahard Diamond Wafer via High-Frequency\nPulsed Local Non-Equilibrium Growth\nREVIEWER COMMENTS\nReviewer #1 (Remarks to the Author):\nI am satisfied with the revision, recommend for publication now.\nResponse: We sincerely thank Reviewer #1 for the positive evaluation of our revised\nmanuscript and for recommending it for publication. We greatly appreciate the reviewer’\ns time and effort in improving the quality of our work.\nReviewer #2 (Remarks to the Author):\nThe authors have thoroughly addressed all reviewer concerns. Taking into account their\nresponses to other referees as well, the scientific integrity of the manuscript has been\nsignificantly improved, and the manuscript is recommended for publication in Nature\nCommunications.\nResponse: We sincerely thank Reviewer #2 for the thoughtful and encouraging\ncomments. We are grateful that the reviewer recognizes our efforts in addressing all\nconcerns and acknowledges the improvement in the scientific integrity of the\nmanuscript. We greatly appreciate the recommendation for publication.\nReviewer #3 (Remarks to the Author):\nThe authors have made substantial efforts to address my initial concerns, and I find their\nwork on synthesizing 5-inch ultrahard diamond wafers to be a notable achievement that\nmerits publication with minor adjustments.\nResponse: We sincerely thank the reviewer for the positive evaluation of our work and\nfor recognizing the significance of our achievement in synthesizing 5-inch ultrahard\ndiamond wafers. We truly appreciate the constructive feedback and are glad that our\nrevisions have addressed the initial concerns. Following the reviewer’s suggestion,\nwe have carefully made the required minor adjustments in the revised manuscript to further\nimprove its clarity and completeness.\nWe believe that these changes strengthen the presentation of our results, and we are\ngrateful for the reviewer’s support and insightful comments, which have helped us to\nimprove the quality of the paper.\nComment 1: However, I suspect the hardness values may still be overestimated. Based\non the cited literature, AFM measurements typically overestimate hardness by about\n20% compared to optical microscopy methods. Even with this correction, yielding\napproximately 180 GPa, this remains impressive for a 5-inch sample and should be\nacknowledged.\nResponse 1: We thank the reviewer for raising this important point regarding the\naccuracy of hardness measurements. Atomic force microscopy (AFM) can directly\ncapture the cross-sectional profile of an indentation and thereby determine the distance\nbetween two vertices on the sample surface reference plane, i.e., the indentation\ndiagonal length. In contrast, when using optical microscopy, small protrusions are often\nformed around the indentation due to plastic deformation, which may cause the\nmeasured diagonal length to appear slightly larger. Since hardness is inversely\nproportional to the square of the diagonal length, the hardness values obtained by AFM\nare therefore higher than those measured by optical microscopy. This issue has been\ndiscussed in detail in Refs. 1 and 2. As shown in the Fig. R11, 2, for the same indentation,\nthe optical method gives an average diagonal length of 9.75 μm, corresponding to a\nnominal Vickers hardness of 191.1 GPa, whereas AFM yields a more accurate average\ndiagonal length of 8.80 μm, corresponding to a hardness of 234.6 GPa. This indicates\nthat the optical method overestimates the indentation size by about 10.8%, which in\nturn underestimates the hardness value by ~18.5% (43.5 GPa).\nIn our work, the indentation morphology was observed and measured using a\nlaser confocal microscopy. This technique essentially belongs to the category of\noptical microscopy, as it is based on optical imaging principles. Therefore, the measured\ndiagonal lengths may be slightly larger than the actual values, which means that the\nreported hardness values may be slightly lower than the actual hardness.\nNevertheless, for the 5-inch diamond sample, here might be some margin of error\nin the hardness measurement. To account for this, we performed multiple indentation\ntests at different sample locations, and the measured hardness values ranged from ~180\nGPa to ~210 GPa. We have added a discussion of this issue in the revised manuscript.\n(Please see Page 8 in the revised manuscript)\n[Figure Redacted]\nFig. R1 AFM calibration of an indentation formed on a nanotwinned diamond sample\nat a load of 9.8 N for hardness and fracture toughness tests. The line profiles shown at\nthe right are colored corresponding to the lines in the AFM image. An optical\nmicroscope photograph is shown as an inset of the AFM image. The mean diagonal is\n8.80 μm from our AFM calibration and 9.75 μm from our optical measurements (cyan\nbox).1\n1. Xu B., Tian Y. J. Ultrahardness: Measurement and Enhancement. The Journal of\nPhysical Chemistry C, 119 (10), 5633-5638 (2015).\n2. Xu B., Tian Y. J. High pressure synthesis of nanotwinned ultrahard materials.\nActa Physica Sinica, 66 (3): 036201 (2017). (In chinese)\nComment 2: Regarding the thermal stability discussion, there is a structural issue with\nthe manuscript organization. The authors discuss stacking fault evolution during\nthermal treatment before they have properly identified and characterized the stacking\nfaults through TEM analysis. This discussion should be repositioned to appear after the\nTEM characterization section where stacking faults are definitively identified and\nanalyzed.\nResponse 2: We appreciate the reviewer’s careful reading and constructive comment.\nWe agree that the discussion of stacking fault evolution during thermal treatment would\nbe clearer and more logically consistent if it follows the TEM characterization section,\nwhere the stacking faults are first identified and analyzed. In the revised manuscript,\nwe have repositioned this part accordingly: the TEM results and stacking fault\ncharacterization are now presented first, followed by the discussion of stacking fault\nevolution during thermal treatment. This adjustment improves the manuscript\norganization and strengthens the logical flow of the results and discussion.\nComment 3: Most concerning are the reported GSFE values, which appear much lower\nthan recent literature including both their cited Carbon paper and a 2023 Physical\nReview B publication (10.1103/PhysRevB.108.014104). This discrepancy requires\nbetter explanation and comparison with current calculations.\nResponse 3: We sincerely thank you for pointing out this issue. In our original\nmanuscript, we only calculated the stacking fault energy of the stable structure (Fig.\nR2c), which resulted in a relatively small value. Following your valuable suggestion,\nwe have now performed generalized stacking fault energy (GSFE) calculations,\naccording to the method described in the referenced paper (Phys. Rev. B 108, 014104,\n2023), to provide a more comprehensive understanding of the energy variation with slip\ndisplacement. The recalculated results are shown in the revised manuscript (see Fig.\n3g). For undoped diamond, the GSFE is obtained as 5.682 J/m2 at 0.5b (Atomic\nstructure see Fig. R2b), which is consistent with the values reported in the literature.\nAfter nitrogen incorporation, the GSFE decreases significantly to 2.226 J/m2, clearly\ndemonstrating that nitrogen doping reduces the stacking fault energy and thus may\nfacilitate fault formation. We have revised the manuscript accordingly, including both\nthe updated GSFE data and a more detailed description of the computational\nmethodology to ensure consistency and clarity.(Please see Page 14 and Page 20 in the\nrevised manuscript)\nFig. R2 Atomic structures at a slip of 0.0b(a), 0.5b(b), and 1.0b(c) for <112> {111}\nslip systems of diamond.\n“Fig. 3g Density functional theory (DFT) calculated GSFE curves on glide-set plane\nof undoped diamond and N-doped diamond”\nDespite these concerns, the work represents a valuable contribution to ultrahard\nmaterials research. The synthesis of large-scale diamond wafers with engineered\nstacking fault networks is a significant technological advancement. With these\nadjustments addressing measurement limitations, manuscript organization, and\nliterature consistency, the manuscript should be acceptable for publication.\nResponse: We sincerely thank the reviewer for the positive evaluation of our work and\nfor recognizing the significance of synthesizing large-scale diamond wafers with\nengineered stacking fault networks. We are also grateful for the constructive\nsuggestions regarding hardness measurement, manuscript organization, and literature\nconsistency. These have all been carefully addressed in the revised version, which we\nbelieve has substantially improved the clarity and overall quality of the manuscript."}, {"round": 3, "reviewer_comments": {"reviewer_3": "the authors have addressed my concerns soundly. I thus recommend the work for publication in NC."}, "author_response": ""}]

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