However, whether this O2 pulse trajectory is within the range of normal has been debated, and the diagnostic performance of the O2 pulse for OCAD in physically fit individuals, in whom V˙O2 may be more likely to plateau, has not been evaluated. Using prospectively collected data from a sports cardiology program, patients were identified who were free of other cardiac disease and underwent clinically‐indicated CPET within 90 days of invasive or computed tomography coronary angiography. The diagnostic performance of quantitative O2 pulse metrics  and qualitative assessment for O2 pulse plateau to predict OCAD was assessed. Among 104 patients , the diagnostic performance for OCAD  was poor for both quantitative and qualitative metrics reflecting an O2 pulse plateau . When O2 pulse parameters were added to the electrocardiogram, the change in AUC was minimal . Those patients without OCAD with a plateau or decline in O2 pulse were fitter than those with linear augmentation  and had a longer exercise ramp time . Overall, a plateau in O2 pulse was not a useful predictor of OCAD in a physically fit population, indicating that the O2 pulse should be integrated with other CPET parameters and may reflect a physiologic limitation of stroke volume and/or O2 extraction during intense exercise.Cardiopulmonary exercise testing (CPET) guidelines recommend analysis of the oxygen (O 2) pulse for a late exercise plateau in evaluation for obstructive coronary artery disease (OCAD). The diagnostic performance of the O2 pulse for OCAD in physically fit individuals, in whom V˙O2 may be more likely to plateau, has not previously been evaluated. We found that a plateau in O2 pulse was not a useful predictor of OCAD in a physically fit population and may reflect a physiologic limitation of stroke volume and/or O2 extraction during exercise rather than a pathologic fall in stroke volume due to ischemia.Cardiopulmonary exercise testing (CPET) guidelines recommend analysis of the oxygen (O Because O2 extraction is typically thought to increase linearly and predictably across the spectrum of exercise capacity,  offers a broad approach to evaluate exertional symptoms, ‐only exercise testing suffers from only moderate sensitivity and specificity for the diagnosis of obstructive coronary artery disease (OCAD),  significant arrhythmia during CPET, (b) respiratory exchange ratio <1.05, (c) poor quality gas exchange measurements, or (d) short ramp duration (<4 min) of the exercise protocol.Participants were eligible for inclusion in this study if they were ≥18 years old and performed CPET on the treadmill or cycle ergometer in the exercise lab of the Massachusetts General Hospital (MGH) Cardiovascular Performance Program (CPP)  within 90 days of clinically‐indicated invasive coronary angiography (ICA) or coronary cardiac computed tomography (CCTA). Rigorous clinical exclusion criteria  through June 1, 2019, CPET results were prospectively collected in a research database. All aspects of this study were approved by the Partner's Human Research Committee . Patients or the public were not involved in the design, or conduct, or reporting, or dissemination plans of the research.2.2Intensity graded, maximal effort‐limited exercise tests were performed with continuous gas exchange on the treadmill  or upright cycle ergometer . The cycle ergometry protocol consisted of 3 min of free‐wheel cycling followed by continual increase in resistance (10–40 watts/min). Treadmill tests began with a 10‐min warm‐up at 3.0 to 7.5 miles/h and 1% grade followed by a progressive increase in incline (0.5% grade increase every 15 s) at a fixed speed. The intensity of the cycle ergometry ramp and the speed of treadmill testing were determined by the exercise physiologist and participant with a goal of an 8–12 min ramp time. All tests proceeded until exhaustion, onset of limiting symptoms, or development of a clinical contraindication to exercise  and analyzed using Breeze Suite software . Peak gas exchange parameters  are reported as the highest average over a period of 30 s during the last minute of effort‐limited exercise. Predicted pV˙O2 was calculated by the Jones equations  was corroborated by observation of minute ventilation, V˙CO2, and HR. Rarely and limited to treadmill tests, when there was a clear decrease in two or more of these parameters prior to when peak work was recorded, peak work time was moved earlier for purposes of O2 pulse analysis. Visual inspection of plotted breath‐by breath V˙O2 was used to qualitatively assess when the slope of V˙O2 through the end of exercise was approximately zero or less than zero, reflecting V˙O2 plateau, and for the purposes of O2 pulse analysis, the timing of peak V˙O2 was considered the beginning of this V˙O2 plateau. A true V˙O2 plateau was defined as greater than 30 s. All O2 pulse curves and slope fits were inspected, and outlier values in the breath‐by‐breath data (>2 standard deviations from the fitted slope) were excluded and slopes were re‐inspected to assure they accurately reflected the shape of O2 pulse curves. Quantitative and qualitative analyses in this study were performed using the breath‐by‐breath data without averaging or smoothing data other than removing outlier values as described above.Gas exchange was measured on a breath‐by‐breath basis using a commercially available metabolic cart (Ultima Cardia OExercise testing was performed with 12‐lead ECG monitoring (Mortara Instrument X12+), and measurement of blood pressures using manual sphygmomanometer  for each patient were created from the start of CPET ramp until peak work and peak V˙O2 (pV˙O2). Both peak work and pV˙O2 were assessed given pV˙O2 may reach a physiologic plateau prior to peak work in a proportion of athletes  the start of ramp until the final 2 min of exercise (Early Exercise Slope), and (b) the final 2 min of exercise to pV˙O2 (Late Exercise Slope). The proportional change in slopes is defined as (Late Exercise Slope – Early Exercise Slope) / Early Exercise Slope. The same assessment was also performed through peak work.All participants underwent an intensity graded, maximal effort‐limited exercise test with continuous gas exchange monitoring on the treadmill or upright cycle ergometer. Individual breath‐by‐breath gas exchange data were extracted and plots of the O2.42 pulse trajectory was performed by two physician reviewers (B.P. and T.C.) who were blinded to patient data including angiography. Tests were categorized into four groups  Flat Throughout: no augmentation of the O2 pulse during the ramp, Category C) Plateau in Late Exercise: initial augmentation of the O2 pulse with flattening during late exercise, Category D) Decline in Late Exercise: initial augmentation of the O2 pulse with down‐sloping during late exercise. If the two reviewers assigned the same test to different categories, a third physician reviewer (M.W.), blinded to the initial categorization, adjudicated the study.Qualitative assessment of O2.5Angiograms were performed as clinically‐indicated at a single center (MGH), and clinical reports were used to assess the presence of CAD. The presence of CAD (“any CAD”) was defined as any luminal irregularities on ICA or CAD‐RADS>0 on CCTA  and compared by the Chi‐square test or Fischer's exact test when n ≤ 5 per category. Univariate predictors for an abnormal O2 pulse trajectory were assessed using logistic regression. The diagnostic performance for OCAD of O2 pulse metrics alone and in combination with the ECG was assessed via receiver operator characteristic curve analysis. Statistical analysis was performed with R: A Language and Environment for Statistical Computing (R Core Team 2021) and SAS (Version 9.4). O2 pulse curve slopes were generated in Graph Pad Prism (version 8.4.1).Descriptive continuous variables are presented as mean (SD) or median [IQR] as appropriate and were compared between groups using a two‐sample unpaired 33.1p < 0.05) but similar body mass index as those on the cycle ergometer. Baseline hemoglobin levels were available in 60% of the cohort and 5 male (Hgb <13) and 1 female athlete (Hgb <12) met criteria for anemia.Among 2,155 patients who were referred for clinically‐indicated CPET, 104 patients met study inclusion criteria  and peak V˙O2 (41.3 ± 10.5 vs. 31.5 ± 11.9 ml/kg/min) than those on the cycle ergometer , 9 athletes had moderate EIAH (SpO2 <93%), and no athletes had severe EIAH (SpO2 <88%) . While peak V˙O2 occurred before peak work in 76% of tests, only 12% of all tests had a true plateau in pV˙O2, defined as greater than 30 s. Results in subsequent sections refer to assessment of the O2 pulse trajectory through pV˙O2 excluding any period of V˙O2 plateau unless otherwise noted.Consistent with lab protocols, which include a longer warm up period on the treadmill, total test duration on the treadmill was longer but included a shorter period of time on the ramp than on the cycle ergometer (Table 3.32 pulse trajectories and the corresponding O2 pulse slopes. Most (64%) tests had normal linear augmentation throughout exercise (Category A). Of the remaining tests, 26% had O2 pulse trajectories that plateaued during late exercise (Category C) with slope decrement of 85 ± 36%, and 7% had O2 pulse trajectories that declined during late exercise (Category D) with slope decrement of 128 ± 19%. Few tests (3%) had O2 pulse that was flat throughout the ramp (Category B). Time on the ramp portion of the test was significantly shorter in these individuals versus patients with the other O2 pulse trajectories , and all of these tests occurred in participants without CAD  changed categories, with more having a late exercise decline , and fewer having linear augmentation .Figure 3.4p < 0.05). PV˙O2 and O2 pulse were similar among groups and remained supra‐normal even in those with multi‐vessel OCAD. The late exercise O2 pulse slope was similar in those with and without CAD, including OCAD. Flattening of the O2 pulse in late exercise relative to early exercise was numerically greater with increasing severity of CAD , but was highly variable within and not significantly different between groups. Most participants with OCAD had normal linear augmentation in the O2 pulse . The proportion of participants with OCAD who had a plateau or decline in the O2 pulse  was similar to the group without any CAD (22%). Overall, there were no significant differences in the distribution of O2 pulse trajectories among the qualitative categories in those with any form of CAD versus those without CAD. Alternate categorization of O2 pulse curve shape using test data through peak work instead of pV˙O2 did not change this result.Table 3.52 pulse parameters had poor predictive capability for OCAD. The AUC for a normal (Category A) versus abnormal  O2 pulse trajectory was 0.52 , and this was not improved by excluding tests with flat O2 pulse trajectory  or classifying only those tests with O2 pulse decline as abnormal . The AUC for late exercise O2 pulse slope was 0.55, with an optimal cut‐off of 0.48 ml/beat/min . Similarly, the AUC for the proportional change in late exercise O2 pulse slope was 0.55, with an optimal cut‐off of 0.55 . When O2 pulse parameters were added to the exercise ECG, the change in AUC was minimal . When the O2 pulse trajectory was evaluated through peak work rather than pV˙O2, results were similar  change when added to ECG).In the evaluation for OCAD, the presence of ischemic ECG changes had an area under the receiver operating curve (AUC) of 0.69 and sensitivity (Se) of 82%, and specificity (Sp) of 57%. All O3.62 pulse parameters for OCAD, alternative predictors for a plateau or decline in the O2 pulse trajectory were explored in Table without OCAD. Participants with a plateau or fall in the O2 pulse reached a higher percent of predicted pV˙O2  and had a longer exercise ramp length (9.5 ± 3.2 vs. 8.0 ± 2.5 mins) as compared to those with normal linear O2 pulse augmentation. There were no other significant differences among O2 pulse trajectory groups in participants without OCAD.Given the poor predictive ability of the O42 pulse trajectory for the diagnosis of OCAD in a physically fit population of patients, free of other relevant forms of cardiac disease, who were referred for clinical CPET. Key findings are summarized as follows. First, a plateau or decline in the O2 pulse trajectory during late exercise, as evaluated by several complementary qualitative and quantitative metrics, was not associated with the presence of OCAD on coronary angiography. Second, there was no significant improvement in diagnostic performance of CPET when O2 pulse metrics were added to exercise ECG. Third, among those without OCAD, an O2 pulse decline or plateau was associated with attainment of higher percentage of predicted pV˙O2 and longer exercise ramp time. Overall, these findings suggest that in a physically fit population, a plateau or decline in O2 pulse during late exercise is not a useful predictor of OCAD.This study sought to evaluate the diagnostic performance of quantitative and qualitative parameters defining O2 pulse trajectory assessment to standard exercise test variables in the diagnostic evaluation of patients with suspected myocardial ischemia , flattening of the O2 pulse and V˙O2 trajectories significantly improved the diagnostic performance of ECG‐only exercise testing for myocardial ischemia against a reference standard based on single photon emission computed tomography (SPECT)  had single vessel disease. Given the small number of individuals with multi‐vessel disease in our study, we cannot exclude the possibility that O2 pulse metrics could have adequate sensitivity for disease affecting multiple coronary territories. Consistent with this finding, Munhoz et al. previously demonstrated that O2 pulse trajectory was impacted by extensive but not mild ischemia on SPECT  of invasive angiograms utilized physiologic testing with FFR or iFR, all raising the possibility of misclassification of CAD severity. However, OCAD was defined by ICA in the vast majority of cases (83%), and we noted that moderate disease was not more common in those with an O2 pulse decline or plateau (Table This study has several limitations that merit discussion. First, our cohort was small and included tests performed on the treadmill and cycle ergometer. We combined the analysis because tests on cycle and treadmill shared similar late exercise Oau Table , lesseni2 pulse metrics to identify a late exercise plateau, as it is recommended by guidelines, was limited by low sensitivity and specificity for OCAD in our physically fit cohort and did not improve the diagnostic performance of ECG. Though O2 pulse metrics may still have a role in the evaluation for ischemia‐induced LV dysfunction in those with average or lower exercise capacity, our results suggest that the use of O2 pulse trajectories in isolation or in combination with the ECG may have limited diagnostic utility in the prediction of OCAD in physically active clinical populations who undergo CPET. Though CPET remains an enormously useful clinical test for the evaluation of cardiopulmonary disease (Balady et al., In conclusion, we found that the application of ONone.


Molecular vanadium oxides, or polyoxovanadates (POVs), have recently emerged as a new class of molecular energy conversion/storage materials, which combine diverse, chemically tunable redox behavior and reversible multielectron storage capabilities. This Review explores current challenges, major breakthroughs, and future opportunities in the use of POVs for energy conversion and storage. The reactivity, advantages, and limitations of POVs are explored, with a focus on their use in lithium and post‐lithium‐ion batteries, redox‐flow batteries, and light‐driven energy conversion. Finally, emerging themes and new research directions are critically assessed to provide inspiration for how this promising materials class can advance research in sustainable energy technologies. This Review critically discusses recent breakthroughs and future challenges in research on polyoxovanadate energy materials. The use of polyoxovanadates in batteries, redox‐flow batteries, light‐driven catalysis, and electrocatalysis is described together with an outlook on emerging themes and areas of future application. The transition from traditional fossil‐fuel‐based systems to carbon‐neutral and more sustainable schemes is underway.12O40]n− and Wells‐Dawson [X2M18O62]m− anions .⋅x H2O . As a consequence of this synthetic approach, the crystal lattice typically contains water, which needs to be thermally removed, as otherwise this would lead to “gassing” (O2 and H2 evolution) under typical LIB operation conditions.6[V10O28] as a LIB cathode material, which led to initial discharge capacities of about 130 mAh g−1 at a current density of 0.2 mA cm−2 and potentials between 2.0 and 4.2 V vs. Li+/Li.−1) but low cycling stability.−1) but high cycling stability.6[V10O28] as the cathode active material in LIBs. By using in situ X‐ray absorption near‐edge spectroscopic (XANES) studies, the authors showed that all ten V5+ ions can be reversibly reduced to V4+ in a potential range of 4–1.75 V vs. Li+/Li.6[V10O28] was reported as an anode material with a reversible capacity of about 280 mAh g−1, an average discharge potential of 0.4 V vs. Na+/Na, and high cycling stability.6[V10O28]⋅16 H2O leads to the formation of two solid‐state oxides, LiVO3 and LiV3O8, even at moderate temperatures of about 120 °C.To date, most studies have been focused on the POV prototype decavanadate, which is typically synthesized under aqueous conditions and isolated as a hydrated salt, that is, M5.72H3.28[PV14O42],7[NiV13O38],7[MnV13O38],7[V15O36(CO3)][7[V15O36(CO3)] as the active anode and cathode material. The system combined battery‐like energy density (125 Wh kg−1) and supercapacitor‐like power density (51.5 kW kg−1 at 100 A g−1) and could serve as a model to bridge these two technologies.In addition to studies using the decavanadate prototype, other common POVs such as KO36(CO3)]67, 68−1 in a voltage range between 1.2 and 3.4 V vs. Li+/Li, at a current density of 50 mA g−1. The viability of this POV stabilization method is supported by related studies by Yoshikawa and co‐workers, who used dicationic N,N‐dimethylbiguanidinium (DBG) to access the crystalline compound [DBG]3[V10O28].−1 in a voltage range between 1.5 and 3.8 V vs. Li+/Li. Further work by Liu and co‐workers showed that the combination of Mg2+ and NH4+ cations in Mg2(NH4)2[V10O28] also results in stable crystal lattices and gives access to LIB cathodes with discharge capacities of about 200 mA h g−1 in the potential range of 1.0 and 3.8 V vs. Li+/Li.To explore the performance of truly molecular POV battery electrodes, Anjass, Streb, and co‐workers proposed the use of supramolecular crystal engineering for POV stabilization.These reports highlight the urgent need for fundamental studies on the reactivity of POVs, with a focus on thermally or chemically harsh conditions as used in many energy technologies. Based on this understanding, new concepts for stabilizing POVs  could open new approaches for the design of materials for the field. In addition, fundamental questions also remain with respect to interfacial reactivity and the role of the nano‐ and microstructuring of POM‐based electrodes,4−1 by the solubility of the vanadate salts.−1.Redox‐flow batteries (RFBs) charge and discharge (highly concentrated) solutions of redox‐active species. The charged solutions can be stored in external tanks, which can be varied in size to allow inexpensive and simple scaling of the battery capacity without altering the electrode area. This results in decoupled power and capacity, which is not possible with other battery technologies and is an advantage for stationary energy‐storage applications.−1,When designing new redox‐active species for RFBs, the two main chemical considerations are 1) how much charge can be stored per liter (volumetric capacity) and 2) how much energy can be stored per charge . Since RFBs are mainly aimed at  stationary energy storage, low materials cost is an important third consideration. Molecular and macromolecular species with multielectron storage capability (such as POMs or redox‐active polymers) can fulfil all three design criteria.78, 79Although most of the POM‐RFB studies have thus far been focused on tungstate POMs, POVs offer some features which make them ideal charge‐storage materials in RFBs. POVs can often be accessed as mixed‐valence species, so that they can be both oxidized and reduced.4[SiW12O40]) anolyte with a POV (Na4.75H4.25[PVV14O42]) catholyte.−2 greater than an all‐vanadium RFB.V14O42]9− was chemically reduced to [PVV10VIV4O42]13−, which could then be oxidized by 4 electrons when charging the RFB catholyte. During charging, the H4[SiW12O40] anolyte was reduced by 2 electrons per cluster. Therefore, the clusters were used in a 2:1 molar ratio (tungstate/vanadate) to ensure that all the redox‐active material could be charged.51V NMR spectroscopy that, upon degradation, the [PV14O42]9− cluster reassembled under operational conditions. This behavior could lead to significantly longer operational lifetimes for POM‐based RFBs compared to other RFB systems.  exhibited solubilities of up to 1.2 m in acetonitrile (containing 0.1 mnBu4NPF6).[V2VIV4O7(OR)12] allows the cluster to be reduced or oxidized by up to 2 electrons. This enabled the development of an RFB, where the same POV is used as the active species in the anolyte and catholyte. This can be advantageous because it prevents long‐term performance loss arising from active species crossover. Matson and co‐workers went on to demonstrate that chemical modification of their Lindqvist cluster can be used to tune the electrochemical performance. For example, tuning the redox potentials by incorporation of heterometals into the cluster framework is possible,In their efforts to increase the energy density, the RFB community has started to explore non‐aqueous RFBs. Switching to organic electrolytes increases the solubility of many redox‐active organic materials, and extends the functional voltage window of RFBs by preventing the electrolysis of water into oxygen and hydrogen.82, 835The ability of POVs to absorb in the visible‐light region makes them well‐suited for photoredox catalysis using sunlight.4O12]4− (4}{V) to give the visible‐light‐absorbing [V5O14]3− (5}{V).IV2VV8O26]4− (IV{V2VV8}). The authors were able to close the catalytic cycle by reoxidation of the decanuclear species using molecular oxygen (slow) or hydrogen peroxide 3)2V12O33]−, which only differ in their internal template, where X is either chloride or bromide. Consequently, the two clusters show almost identical light absorption. The authors examined the photoredox activity of both clusters under identical conditions, using the photooxidative degradation of the model pollutant patent blue V (PBV) as a test reaction. This comparative study showed that the bromide‐templated system exhibited significantly faster PBV degradation kinetics compared with the chloride species. By using ultrafast photophysical analyses together with theoretical computations, it was proposed that, upon photoexcitation, the heavier bromide template enables a more efficient singlet–triplet transition because of the enhanced spin–orbit coupling .2 from the gas to the liquid phase can severely limit the catalytic performance of POVs.Building on these studies of metal‐functionalized POV photooxidation catalysts, Guldi, Streb, and co‐workers demonstrated that modification of the internal cluster template anion can also be used to modulate the photoreactivity. The authors explored the model compounds [4V4O17(OAc)3]3− (4V4}{Mn) as a functional model of the oxygen‐evolving complex [CaMn4O5], which enables water oxidation in the natural photosystem II.3]2+  and the terminal oxidant persulfate, the catalyst shows water oxidation under irradiation with visible light.−1 were achieved with this system.The fields of light‐driven POV catalysis and POV metal functionalization have recently been combined to give access to bio‐inspired POV‐based water oxidation catalysts. In one example, Streb and co‐workers reported the manganese‐functionalized POV 6−, [PV14O42]9−, and [V6O7(OR)12]. This is a stark contrast to the fast‐paced recent progress in the design and functionalization of new POV architectures—many of which have not been studied for their possible use in energy technologies.[6O7(OR)12] to tune the redox reactivity and electron‐storage capability.[12O32Cl]5− is a facile route to tune redox behavior.2+ ions into the dodecavanadate cluster leads to a significant enhancement of the cluster redox activity, thus allowing the fabrication of LIB cathodes with improved energy density. The authors attributed this surprising finding to the electrostatic and structural stabilization of the reduced cluster by Ca2+ ions.One major challenge in future POV energy research is a comprehensive analysis of POV reactivity and stability. This requires POV experts to collaborate closely with materials scientists, electrochemists, and photochemists to gain a deeper understanding of POV structure–function relationships. This could be used to develop new POVs with tunable reactivity and identify new POV structures as targets of future research. To date, most POV energy applications have been focused on a few POV prototypes, most notably  as hole scavengers on semiconducting CdSe light absorbers and demonstrated that this system shows enhanced light‐driven hydrogen evolution. The authors assign this observation to the ability of the POV to act as a hole scavenger, which facilitates proton reduction on the semiconductor surface.An alternative approach towards higher visible‐light photoactivity is the coupling of POVs with metal complexes or organic photosensitizers,10O32]4− was used as a photocatalyst to trigger the abstraction of a C−H hydrogen atom, which enabled the arylationIn the field of C−H activation, MacMillan and co‐workers recently reported a ground‐breaking study in which the tungstate archetype  as an electrode material for a supercapacitor,−1 and a power density of 312 W kg−1, based on double‐layer and pseudocapacitance. The study lays the foundations to expand POV research into the supercapacitor domain.In addition, the use of POVs in supercapacitors has thus far not been explored in detail. Supercapacitors can store energy through a combination of electrical double layer capacitance and pseudocapacitance ,2, O2, CO2, N2). In one recent example, Streb and co‐workers demonstrated that the molecular‐water oxidation catalyst 4V4}{Mn (see Section 4) also shows sustained oxygen evolution under electrochemical conditions;4− to bind elemental Br2 in the central cavity, thereby leading to polarization of the Br−Br bond and significant selectivity changes in organic bromination reactions.2 from air and store it as an internal CO32− template, thereby resulting in the species [H8VIV15O36(CO3)]6−.In the field of small‐molecule activation, Kikukawa et al. have recently broken new ground by using the bowl‐shaped dodecavanadate [V7POVs have proven to be versatile redox‐active materials that can be used to overcome challenges in many sustainable energy applications, from batteries to light‐driven catalysis. Despite this, only a small fraction of the known POV cluster types have been investigated as components of energy conversion or storage systems. Furthermore, more extensive investigation of the physical properties of existing species is needed to use them effectively in energy applications, such as, determining trends in redox behavior, photophysical properties, thermal stability, and solubility. Expanding the available knowledge of POV properties will allow the field to move beyond the model POVs described here and allow POVs to be designed to overcome current limitations. By addressing the issues we have highlighted, we believe that POVs can become key components in the development of future sustainable energy storage systems.The authors declare no conflict of interest.Montaha Anjass is an independent group leader and Margarete‐von‐Wrangell‐Fellow at Ulm University and Helmholtz Institute Ulm. She received her B.Sc. in Chemistry from Birzeit University, Palestine, and her M.Sc. in Advanced Materials from Ulm University, Germany. In 2019, she completed her PhD  at Ulm University. Her current research interests are advanced battery materials based on redox‐active molecular metal oxides and electrically conductive organic polymers, and their use in (post‐)lithium batteries.Grace A. Lowe is a postdoctoral researcher in the Streb group at Ulm University. She completed her PhD at the Centre for Doctoral Training in Sustainable Chemistry at Nottingham University in 2019 under the supervision of Dr. D. A. Walsh and Dr. G. N. Newton. Her current research interests are redox‐flow batteries and the electrochemical characterization of inorganic materials, mainly, but not limited to, POMs.Carsten Streb is Professor of Inorganic Chemistry at Ulm University and group leader at the Helmholtz‐Institute Ulm. His love for POMs started during his PhD, where he worked with Lee Cronin at the University of Glasgow. His current research is focused on designing POM‐based functional materials and composites to address global chemical challenges, including energy conversion/storage, water purification, and public health.


In situ XRP allows to follow textural changes of a dried gel Ni/Al2O3 sample as a function of temperature during calcination, activation and CO2 methanation reaction. Complementary PXCT studies on dried gel particles of Ni/Al2O3 and Al2O3 provide quantitative information on pore structure, size distribution, and shape with 3D spatial resolution approaching 50 nm, while identical particles are imaged ex situ before and after calcination. The X‐ray imaging results are correlated with N2‐sorption, Hg porosimetry and He pycnometry pore characterization. Hard X‐ray nanotomography is highlighted to derive fine structural details including tortuosity, branching nodes, and closed pores, which are relevant in understanding transport phenomena during chemical reactions. XRP and PXCT are enabling technologies to understand complex synthesis pathways of porous materials.The synthesis of hierarchically porous materials usually requires complex experimental procedures, often based around extensive trial and error approaches. One common synthesis strategy is the sol–gel method, although the relation between synthesis parameters, material structure and function has not been widely explored. Here, in situ 2D hard X‐ray ptychography (XRP) and 3D ptychographic X‐ray computed tomography (PXCT) are applied to monitor the development of hierarchical porosity in Ni/Al 2O3 catalysts are investigated with sub‐100 nm resolution.Hard X‐ray ptychography in in situ 2D and ex situ 3D modes is shown as an emerging technology to study textural and structural changes of porous materials during thermal treatment. The porosity changes from the dried gel to calcined state of hierarchical porous (Ni/)Al Diffusion limitations of reactants or products are highly relevant both in normal operating conditions and following changes induced by catalyst deactivation. Therefore, strategies have to be found for a knowledge‐based design of porous catalysts with optimized mass transport properties and improved catalyst performance. One promising approach is to design catalysts with hierarchical rather than uniform pore structures, generally including some combination of solids with micro‐ (<2 nm), meso‐ (2−50 nm), or macroporous (>50 nm) pore size distribution. Here, we focus on meso‐/macroporous systems with the predominant transport processes of Knudsen diffusion in mesopores and molecular diffusion or convective flow in macropores. Introducing macropores into a uniform mesoporous system can provide higher mass transfer velocity by promoting convective flow over diffusion.Hierarchical structures are omnipresent in our environment and are particularly found in connection with transport processes, for example, the evolved human respiratory system for oxygen delivery in the lungs, or modern designed road networks for managing traffic flow.1, 2 Common synthesis strategies can be broadly divided into templating, post processing (often leaching), or self‐assembly techniques, each allowing for control over porosity, particle shape and being suited for certain material classes. Modified sol–gel synthesis is a self‐assembly strategy to obtain powders or shaped particles with meso‐/macroporous structure in the form of spheres, membranes, or monoliths. The sol–gel synthesis can be divided into five steps, each of which can influence the final texture and structure of the material: i) sol formation by hydrolysis and partial condensation, ii) gel formation via polycondensation, iii) aging of the gel, iv) drying of the gel, and v) stabilization, specifically calcination. A full understanding of sol–gel synthesis is challenging due to the many steps and parameters involved at each stage. Aside from introducing structure‐directing or phase separation agents , the final porosity can also be strongly determined by the solvent and its removal during drying or calcination. The final calcination step can induce numerous changes to porosity, pore sizes, crystallization, sintering, partial, or even complete structural collapse. Obtaining greater understanding and control over these effects is highly desirable in the context of rational material design.The last two decades have brought significant progress in synthesis and design of meso‐/macroporous materials with potential applications in heterogeneous catalysis or adsorption.19 Com For meso‐/macroporous catalysts a combination of these techniques is necessary due to the different length scales probed, with Hg porosimetry normally used to measure macro‐ and mesopores down to 4 nm, and Ar‐ or N2‐sorption applied for smaller mesopores. However, both methods have limited use in studying hierarchical pore structures, especially to retrieve more complex pore network information like connectivity or shape of individual pores or pore systems. Furthermore, gas sorption studies typically require thermal outgassing during sample preparation, which limits their application in studying dried gel samples and excludes their use in studying late stages of synthesis and some post‐synthesis treatment steps. In addition, Hg porosimetry is considered an invasive method, as probably not all Hg can be removed after the experiment and especially for fragile materials the sample may be damaged by high pressure intrusion of Hg. The application of conventional methods to study pore structure evolution especially during calcination is strictly limited, while in situ experiments are hardly feasible.Following synthesis, a further challenge is the accurate characterization of hierarchical porosity. Schlumberger et al. and Cychosz et al. reviewed gas sorption and Hg porosimetry as routine methods for analyzing hierarchical pore structures.23 For As with conventional methods, meso‐/macroporous catalysts require a combination of imaging techniques to address all relevant length scales (nm–cm). Examples include combining electron tomography for nanoscale analysis, with focused ion beam‐scanning electron microscopy or X‐ray tomography for larger length scales. Among reported X‐ray imaging techniques, hard X‐ray ptychography (XRP) and its 3D counterpart ptychographic X‐ray computed tomography (PXCT) are particularly promising. XRP and PXCT are relatively modern scanning coherent diffraction imaging techniques delivering high spatial resolutions  limited neither by focusing optics (beam size) or by the scanning step size, and applicable to extended samples volumes . XRP and PXCT are capable of quantifying the local electron density (Ne) of the sample in 2D and 3D space, respectively. This in turn provides sensitivity to mass density, which may in principle be used to retrieve chemical information. The application of PXCT in catalysis research is still at a relatively early stage, although the potential of the method is clearly shown in several studies of textural/pore structures or chemical heterogeneity. Recently, PXCT studies on PTFE membranes could derive porosity and tortuosity information, which was then used to calculate diffusion coefficients and permeability.Alternatively, advanced imaging techniques based on X‐ray and electron tomography offer excellent opportunities to study textural and structural properties of hierarchically porous catalysts in a 3D spatially‐resolved manner.29 As 2O3 catalysts during calcination  over a purely mesoporous catalyst (86.5%). Theoretical calculations of the effectiveness factor also suggest a superior performance of the meso‐/macroporous catalyst for methanation of CO2. In both cases, improved performance can be attributed to the hierarchical pore system, stimulating the present aim to investigate how hierarchical porosity can be accurately measured and quantified in complex solids. In situ XRP studies were performed in a nanoreactor setup, ensuring defined gas and temperature conditions during calcination of a dried gel Ni/Al2O3 sample. Complementary PXCT studies of identical dried gel particles of Ni/Al2O3 and Al2O3 before and after calcination were used to quantify the pore structure in 3D space. The X‐ray imaging results were compared with N2‐sorption, Hg porosimetry, and He pycnometry for bulk measurement of meso‐/macroporosity and solid skeletal density. A robust strategy is therefore presented based on XRP to study textural and structural changes at nanoscale resolution, and PXCT to quantitatively probe hierarchical porous features, in both cases producing data largely unavailable to contemporary methods. The methodology shown is broadly relevant for synthesis and characterization of materials with tailored pore structures.In this study, we present a combination of in situ 2D XRP and ex situ 3D PXCT to study the pore structure evolution of meso‐/macroporous Ni/Aln Figure 1. The h22.12O3 and Al2O3 catalysts was adapted from previous studies, reporting bimodal meso‐ and macroporous structures by sol–gel synthesis with polymerization induced phase separation. Briefly, chloride salts of Al and Ni were homogeneously mixed with poly(ethylene oxide) (PEO) in an ethanolic solution. Polymerization and phase separation were induced by addition of propylene oxide (PO) to obtain the wet gels, which were then dried and subsequently calcined. The dried gel samples discussed in this study (Ni/Al2O3‐dr and Al2O3‐dr) were separated from the main synthesis batch after drying, and therefore represent the catalyst state without subsequent treatment steps such as calcination, activation, or catalytic testing (see Supporting Information).Preparation of the hierarchically porous Ni/Alparation.54 BriFigure Elemental analysis  decreased after calcination from 1.5 to 1.0 µm in Ni/Al2O3, and from 1.2 to 1.0 µm in Al2O3∗. It should be noted that Hg porosimetry studies to compare the dried gel and calcined state of Al2O3 were carried out on a different sample batch than for the X‐ray imaging studies (indicated by an asterisk “*”). This was done due to an insufficient amount of sample remaining from the X‐ray imaging batch to allow both Hg porosimetry and N2 sorption, rather only Hg‐porosimetry on the X‐ray imaging batch was performed . To compensate for this, a new batch of dried gel sample (Al2O3‐dr*) was made by the same exact preparation procedure and examined by Hg porosimetry before and after calcination . This might influence the absolute pore width values, as wP, M, Hg for the same Al2O3‐calc batch as the X‐ray imaging studies was 0.82 µm, compared to 1.0 µm for the new batch . Furthermore, small mesopore contributions (<10 nm) were found for Al2O3‐dr at wP, M, Hg close to the resolution limit of Hg porosimetry. The Brunauer‐Emmett‐Teller (BET) surface area increased from 5 m2 g−1 (Ni/Al2O3‐dr) to 99 m2 g−1  after calcination, while a decrease from 128 m2 g−1 (Al2O3‐dr) to 91 m2 g−1  was observed.The mode of the macropore width  Wildfire Nano‐Chip  by gluing using ion‐beam‐induced deposition (IBID) of a Pt precursor inside the FIB system . In situ XRP studies were carried out employing three different gas conditions: i) calcination (20% O2/He), ii) activation (25% H2/He), and iii) reaction (20% H2/5% CO2/He) with a constant total gas flow of 2 mL min−1  . Despite measuring changes of the sample width, deriving further quantitative information about the pore structure is not feasible based on 2D projection data. Further PXCT experiments were therefore performed. In principle, any structural changes of the sample leading to changes in electron density can be investigated by XRP. For Ni/Al2O3 this could be exemplarily the reduction of NiO to Ni during catalyst activation. However, the achieved 2D resolution of ≈67 nm is not sufficient to directly resolve any large Ni particles which may be formed by sintering or aggregation during or after the activation step. Generally, changes in electron density can also be observed even if single individual features such as Ni particles cannot be directly resolved, as shown previously for coke deposition on a Ni/Al2O3 catalyst. In the present case, a quantitative analysis of the electron density is limited due to the tilting of the sample upon thermal treatment, which would require to account for changes of the sample thickness that cannot be directly obtained from only 2D projections. In summary, the ability to resolve structural changes in 2D can be challenging on extended 3D samples and is limited by the achieved 2D resolution and potential textural changes.The in situ XRP studies show the potential of XRP for studying extended sample volumes  with sub‐100 nm resolution under well defined temperature and gas conditions. Comparable results from SEM or FIB‐SEM imaging are not feasible, due to insensitivity to the bulk sample or destruction of the sample, respectively. Likewise the sample volume probed here using hard X‐ray imaging is orders of magnitude larger than could be studied with transmission electron microscopy (TEM), due to electron beam attenuation. While in situ XRP studies in catalysis science were already reported before for sintering experiments on model systems,58 we 2.32O3 and Al2O3 samples, and to recover quantitative Ne values for the hierarchically porous solids. Suitable particles of Ni/Al2O3‐dr and Al2O3‐dr were mounted on Cu tomography pins of the OMNY design by gluing using IBID of a Pt precursor gas inside the FIB‐SEM. The mounted particles had a diameter of ≈50 µm and ≈15 µm for Ni/Al2O3‐dr and Al2O3‐dr, respectively, as shown by SEM  of 91.2 and 55.6 nm, respectively  as described in refs. [Ne tomograms are shown in Figure Ne  of the measured particle volumes (see Supporting Information  for details). The porosity analysis results are summarized in Table M for Ni/Al2O3‐dr showed only a minor increase after calcination to Ni/Al2O3‐calc, while both states showed comparable macroporosity as in a previous PXCT study of the same material after activation (ϵM = 46.3%). For Al2O3‐dr the ϵM slightly decreased after calcination to Al2O3‐calc. For both catalyst types, the Ne slices of similar particle regions showed significant volume shrinkage comparable to the in situ XRP experiments . For Al2O3, the decrease of d¯eq was more pronounced after calcination, also clearly visible in the deq distribution . The retrieved pore network models consist of connected spatial graphs, constructed from pore nodes connected to each other via pore segments. The pore network models for Ni/Al2O3‐dr, Ni/Al2O3‐calc, Al2O3‐dr, and Al2O3‐calc are plotted in , segment volume, segment tortuosity (τ), and segment length  can be retrieved. For each node, the location in the particle, coordination number, and thus amount of terminating and branching nodes, can be analyzed. The location of nodes and segments is defined here by the distance to the gravimetric center of the particle volume (dcenter), while the segment coordinates are approximated as the center point of the connected nodes. Statistical analysis results of the pore network models are shown in Figure 2O3 and Figure 2O3, respectively. Figure rpore distribution of each segment relative to dcenter for Ni/Al2O3 and Al2O3, with scatter sizes indicating normalized relative segment volume . For both samples, calcination generally led to a decrease in the macropore radii (rpore) and appeared to indicate particle volume shrinkage, with a higher density of segments observed at lower dcenter values after calcination. The significance of the segment rpore distribution was tested by performing a kernel density estimation (KDE) of the volume weighted 2D histograms  calculated from the spatial graphs therefore also decreased after calcination for both catalysts as expected  and coordination number of the nodes (CN) can be derived  is also reported  was 1.1 for all four samples. As the segments are short compared to the sample volume investigated by PXCT, τ¯ should not be mistaken for macroscopic tortuosity, which is typically employed for diffusion calculations. However, τ can be used to search for specific regions in the pore network that might form a bottleneck for mass transport in such highly connected pore structures. Therefore, CN¯ and τ¯ were analyzed regarding their dependence on dcenter as shown in Figures CN¯, CN¯br, and τ¯ confirming a homogeneous macropore structure. Furthermore, this quantitatively confirms the observations made during the in situ XRP experiments, indicating no obvious pore opening, closing, or disruption of the macropore network.In addition to quantification of porosity characteristics, further information is readily available from tomographic image analysis, which cannot be directly retrieved from Ned Table . Here thed Table , which iulations.63 How2.5Ne can be obtained, which is indicative of the material composition at each sampling point regardless of the exact spatial resolution. The overall Ne histogram of the whole particle volumes measured by PXCT is shown in Figure 2O3 and Al2O3, respectively. Three main features were observed and labeled, corresponding to the image segmentation procedure discussed previously: Ne ≈ 0 e− Å−3 as air/macropores (“pores” label), Ne ≈ 0.5 to 0.55 e− Å−3 as nanoporous material , and Ne > 0.9 e− Å−3 labeled as “contamination” from sample preparation. Already in the overall Ne distribution  observed by He pycnometry after calcination of the dried gel samples (Table N¯e) of the whole particle volume for Ni/Al2O3 and Al2O3, respectively. Slice dependent analysis allows for easy identification and exclusion of regions with strong contamination from further analysis. For the segmented “material” label, the overall Ne distribution is shown in Figure N¯e in Figure 2O3 and Al2O3, respectively. For Ni/Al2O3, a homogeneous and significant shift of N¯e is found for the selected slice range  for detailed analysis of the majority of macropores. On the other hand, direct quantification of mesopore features below the spatial resolution limit was neither feasible nor statistically meaningful. However, an advantage of PXCT compared to other tomographic imaging methods is that quantitative information about the local n Figure , and inds Figure  of the we Figure . In the d Figure , while t3 Figure  d3 probaNe can be converted into observed mass density (ρ), in order to indirectly measure porosity even below the spatial resolution limit of the measurement. A detailed description of all porosity calculations and derivation is given in the Supporting Information. The porosity of the “material” label (ϵmat) can be derived by comparing the observed (ρ) and theoretical mass density (ρtheo). Subsequently, ϵmat was used to calculate the mesoporosity (ϵm) and total porosity (ϵtot) of the samples, with ϵtot as the sum of ϵM and ϵm. The theoretical density of the structures can be obtained from a known crystal structure or approximated experimentally as ρskel by He pycnometry . It is therefore important to note that only for the calcined states do the above porosity values correspond to true porosity. Since ϵmat, ϵm, and ϵtot are calculated based on sub‐resolution information, that is, local Ne of the material label, therefore the porosity values of the dried gels are influenced by organic and solvent residues, potential intermediate Al2O3 and Ni/Al2O3 structures or blocked pores which cannot be directly resolved. In this respect, mesoporosity values are described as an “apparent” ϵm for the dried gels, which can only be carefully compared to conventional pore characterization methods. In more detail, the “apparent” ϵm for the dried gels from PXCT can hardly resolve or distinguish between accessible or inaccessible porosity, as it is a sub‐resolution information. This is contrary to conventional methods, that is, Hg porosimetry and N2 sorption, which indeed can only resolve accessible pores to the respective probe species (Hg or N2). Therefore, ϵm from PXCT experiments and conventional methods can only be strictly compared for the calcined samples. In the calcined samples, ϵm is not influenced anymore by solvent or organic residues, which might fill the mesopores in the dried gel state making them inaccessible for conventional methods. Nevertheless, this information provides valuable and potentially quantitative insights into sample changes occurring during calcination. On calcination of Ni/Al2O3‐dr to form Ni/Al2O3‐calc, ϵm decreased from 29.4% to 23.4% and ϵtot from 73.3% to 67.9%. A similar trend was also observed for Al2O3‐dr to Al2O3‐calc, where ϵm decreased from 34.3% to 30.1% and ϵtot from 69.2% to 63.7%. Thus, during calcination a significant densification of the nanoporous solid occurs for both systems. ϵm of Ni/Al2O3‐calc is well in line with the value of 24.1% previously reported by electron tomography, where direct imaging of mesopores is feasible due to much higher spatial resolution than PXCT. It is important to note however, that electron tomography is restricted to samples with several orders of magnitude smaller volume than can be measured with hard X‐ray tomography. Particularly for hierarchically porous or complex catalysts, X‐ray microscopy potentially provides a more representative view of the sample structure. The derived ϵm is also in agreement with previous Hg porosimetry combined with He pycnometry. However, conventional porosimetry techniques probe only the bulk sample and cannot provide spatially‐resolved data comparable to that produced by tomography. The agreement with electron tomography and Hg porosimetry therefore validates the quantitative Ne based porosity analysis enabled here by PXCT, even if the resolution is currently not high enough to directly visualize mesopores in the sample.With prior knowledge about the chemical composition, the observed ed Table . An exac2.6 In the current study, the presence of macropores in Al2O3‐dr and Ni/Al2O3‐dr was found by Hg porosimetry, which was shown by PXCT, XRP, and SEM imaging to be an ordered and connected macropore network. While the resolution of PXCT obtained here was not sufficient to study mesopores directly, Ne analysis revealed an apparent ϵm of about 29% and 34% for Al2O3‐dr and Ni/Al2O3‐dr, respectively. From these results it may be concluded that the mesopore system is probably already preorganized in the dried gels, but filled with solvent or polymer species, while exact knowledge about the precursors oxide structure is missing. During calcination, PXCT and in situ XRP revealed an extensive overall volume shrinkage accompanied with a densification of the nanoporous solid, which is also in line with observations during calcination in the laboratory synthesis. While in situ XRP could monitor the changes occurring in real time under controlled gas and temperature conditions, PXCT was required to quantify these changes in terms of pore structure. Based on PXCT, the sample macroporosity was almost constant during calcination, while the mesoporosity and total porosity decreased for both systems  and their influence on the final pore structure is challenging to unravel. In previous work, Tokudome et al. proposed that mesopores were generated during calcination by crystallization of alumina, as they were not detected by Hg porosimetry in the dried gels.54 In  On the other hand, a combination with ex situ PXCT allows to directly retrieve quantitative 3D spatially‐resolved textural and structural information about different states of the synthesis. Additionally, combining X‐ray ptychography with other contrast modes, like X‐ray fluorescence, X‐ray diffraction or X‐ray absorption spectroscopy can lead to further complementary insights into the studied material. Comprehensively, in situ XRP and PXCT provide the opportunity to develop fundamental understanding of complex synthesis procedures, which can be generally applied to understand the synthesis of functional materials, specifically hierarchically porous structures in the present case.The experimental strategy shown here can be generally employed to study textural and structural changes of materials with high spatial resolution, for example during thermal treatment as shown here. In situ XRP can be currently performed in 2D experiments, while a future extension into 3D studies is feasible. This allows to actually follow the evolution of porous materials during their synthesis and to elucidate the influence of different process steps that are typically required. The method is not only limited to the sol–gel strategy discussed in this work, but can be readily extended to templating, casting or other self‐assembly strategies which are common in literature.19 On 32O3 and Al2O3 materials during calcination from dried gels. The calcined materials exhibit a hierarchical structure of meso‐ and macropores, which can improve the mass transport properties and overall performance in catalytic applications. However, a detailed understanding of the complex synthesis procedure to tailor such pore structures is elusive and cannot be easily determined by conventional pore characterization techniques, that is, gas sorption and Hg porosimetry. Here, hard X‐ray ptychography and ptychographic tomography are highlighted as enabling techniques to study textural and structural changes during calcination, achieving high spatial resolution down to 50 nm on extended sample volumes up to 50 µm in diameter. From the dried gel to the calcined samples the macropore diameter decreases significantly for both systems, Ni/Al2O3 and Al2O3. In addition to direct imaging covering almost the entire macropore length scale, PXCT can quantify and track the evolution of local electron density in the nanoporous solid. The electron density increases upon calcination showing a densification of the nanoporous solid, which leads to a decrease of the apparent mesoporosity during calcination. However, as the structure of the dried gels is not known exactly, further targeted in situ studies such as X‐ray total scattering or X‐ray absorption spectroscopy during calcination could be used to further differentiate the influence of the precursor oxide structure and polymer, solvent or H2O. With continuing progress in the development of fourth generation or diffraction‐limited synchrotron sources, PXCT is expected to routinely reach sub 10 nm resolution on similar sample volumes as in the present study, and to permit more rapid acquisition of tomograms than shown here. This would enable to directly study both larger mesopores and macropores over extended sample volumes, potentially also in a high throughput manner. Hard X‐ray ptychography is emerging as a key technique for understanding textural and structural changes in the synthesis of hierarchically porous materials, with applications in heterogeneous catalysis, adsorbents, membranes and numerous other fields.This study details the pore structure evolution of sol–gel synthesized Ni/Alications.52 HowThe authors declare no conflict of interest.Supporting InformationClick here for additional data file.Supplemental Movie 1Click here for additional data file.


The validity of the predicted critical strains crDRXε was then experimentally verified. Fine dynamically recrystallized grains, which formed at the boundaries of the original austenitic grains, were detected in the microstructure of additionally deformed specimens from low-carbon investigated steels. Furthermore, equations describing with perfect accuracy a simple linear dependence of the critical strain crDRXε on peak strain pε were derived for all investigated steels. The determined hot deformation activation energy Q decreased with increasing carbon content  in all investigated steels. A logarithmic equation described this dependency with reasonable accuracy. Individual flow stress curves of the investigated steels were mathematically described using the Cingara and McQueen model, while the predicted flow stresses showed excellent accuracy, especially in the strains ranging from 0 to pε.The hot deformation behavior of selected non-alloyed carbon steels was investigated by isothermal continuous uniaxial compression tests. Based on the analysis of experimentally determined flow stress curves, material constants suitable for predicting peak flow stress  The accuracy of the predicted values pσ and pε was simply evaluated using the correlation coefficient R (−), the relative calculation error Δ (%) and its mean values Δmean (%)—see The determined material constants nd W see  can be upσ is confirmed by high values of correlation coefficients R and very low values of mean calculation error Δmean given in pε the accuracy of back-calculation was lower but entirely sufficient—see pσ a pε depending on the Zener–Hollomon parameter of all investigated steels is shown in pσ and pε was correct, respectively that Equations (5) and (6)  necessary to induce dynamic recrystallization, it is necessary to derive a relationship that will allow the prediction of strain hardening exponent C (see Equation (16)). The value of the strain hardening exponent C is determined for each measured flow stress curve  as the line slope (dependence (15)), considering the zero value of the intersection with the vertical axis. An example of determining the constant C by linear regression of the measured data for one combination of temperature and strain rate (1000 °C and 20 s−1) for steel B is shown in To predict the critical strain C is possible based on its functional dependence on the Zener–Hollomon parameter Z (4), which includes the interaction effects of strain rate and temperature. The dependence of exponent C on parameter Z has a power character for a constant value of a strain rate and a variable deformation temperature [1C (s) and 2C (−) are the material parameters, which are determined by regression separately for each strain rate y-axis), by converting relation (19) into a linear form—see C will, therefore, have the general form:The prediction of the strain hardening exponent perature :(19)C=C1m—see ε˙ —see Figua (s2), b (−), c (s) and d (−) intended for the prediction of the strain hardening exponent C of all investigated steels (according to relation (20)) are given in The values of material constants crDRXε strains necessary to induce dynamic recrystallization of all investigated non-alloyed carbon steels for the given thermomechanical forming conditions. The critical flow stress crDRXσ (MPa), corresponding to the onset of dynamic recrystallization, can then be relatively easily determined using a modified Cingara and McQueen relationship (14), in which the strain ε is replaced by the critical strain required to induce dynamic recrystallization crDRXε:Using Equations (4), (6), (16) and (20) and using the material constants listed in crDRXσ and crDRXε depending on the Zener–Hollomon parameter of all investigated steels is shown in The course of the determined values crDRXε, additional isothermal uniaxial compression tests were performed on steels A, B and D at selected temperature and strain rate combinations. These additional experiments are aimed to detect dynamically recrystallized grains in the microstructure of the investigated steels. The heating mode of individual steels, or cylindrical specimens with a diameter of 10 mm and a height of 15 mm, corresponded to the description given in crDRXε, while the critical strain necessary to induce dynamic recrystallization of crDRXε was determined according to relation (16), resp. according to the procedure described in To verify the validity of the determined critical strains necessary to induce dynamic recrystallization The deformed and quenched specimens were then tempered at 300 °C. The specimens were cut in the middle of their diameter (in the direction of their height) into two halves for metallographic analysis. The specimens prepared in this way were then etched with Alkilo with the addition of HCl to highlight the boundaries of the original austenitic grains. The microstructure in the middle of the height of these specimens was documented by SEM analysis on a JEOL JSM-6490LV microscope under BES imaging . Documentation of the microstructure of the examined specimens from steels A and B is given in crDRXε, which are necessary to induce dynamic recrystallization of steels A and B.Using SEM analysis of uniaxial compression tested specimens, fine dynamically recrystallized grains were detected in the microstructure of the steels A and B in all investigated cases—see Unfortunately, in the case of the steel D, the boundaries of the original austenitic grains could not be clearly etched, not even with the use of other etchants. Signs of the boundaries of the original austenitic grains could only be seen in specimen D3 see , but it pσ and peak strain pε of the investigated steels for the given thermomechanical forming conditions can be predicted—see pσ and the peak strain pε of the investigated steels increases with the increasing value of the Zener–Hollomon parameter—see pε, the situation is more complicated. The low carbon steel A showed the lowest peak strain pε of all investigated steels—see Z (above 1 × 1014 s−1), smaller strains are needed to achieve peak flow stress than in the case of the low carbon steel B—see pε in the whole range of the parameter Z are slightly lower than in medium carbon steel C. At values of the Z parameter above 4 × 1012 s−1, the predicted peak strain pε of the steel D is also lower than in the low carbon steel B—see Using Equations (4) and (6) and using the material constants listed in crDRXσ and critical strains crDRXε necessary to induce dynamic recrystallization, depending on the Zener–Hollomon parameter Z, is shown for all investigated steels in crDRXσ—see pσ. Influence of carbon content or carbon equivalent to the value of the critical strain crDRXε was, except for the high carbon steel D, similar to the critical flow stress crDRXσ. For a similar size of the Zener–Hollomon parameter, smaller critical strains crDRXε are required to induce dynamic recrystallization of the high carbon steel D than in the case of the medium carbon steel C—see The comparison of the determined critical flow stresses crDRXε as well as the values of the peak strain pε are very similar .pσ or crDRXσ and the strain pε or crDRXε increase in all investigated steels, with the increasing value of the parameter Z. This confirms the well-known fact that it is necessary to apply larger strains to induce dynamic recrystallization of the material at high values of the parameter Z [The graphical dependencies shown in ameter Z ,5,40,57,ameter Z .pε and critical strain crDRXε for the given thermomechanical forming conditions of all investigated steels. It is clear from pε and crDRXε increases as the deformation temperature decreases and the strain rate increases. At high deformation temperatures and low strain rates, the difference between pε and crDRXε is small (approx. 0.04 to 0.07). However, in the case of a combination of low deformation temperatures and high strain rates, the difference between pε and crDRXε reaches values up to 0.46, which is a significant fact, especially from the point of view of operating conditions. However, the ratio between crDRXε and pε is very similar for all combinations of thermomechanical forming conditions of all investigated steels—i.e., ranging from 38% to 48%.crDRXε required to induce dynamic recrystallization, which is given in crDRXε can be related to the predicted peak strain pε—see crDRXε depends linearly on the peak strain pε. Thus, for the investigated non-alloyed carbon steels, it is possible, very simply and at the same time with very good accuracy, to determine the value of the critical strain crDRXε only via relation to the value of the peak strain pε:The procedure for determining the critical strains R2 = 0.9955; for Equation (23) R2 = 0.9975; for Equation (24) R2 = 0.9922 and for Equation (25) R2 = 0.9912.The very good accuracy of the above listed equations is documented by the high values of the coefficients of determination: for Equation (22) y-axis is equal to 0, and the slope of the above linear dependencies of the critical strain crDRXε on the peak strain pε of all investigated steels are very similar (see Equations (22)–(25)), it is possible to compile a relationship that would describe this dependence –(25), the intersection with the ence see a comprehcrDRXσ in dependence on the peak flow stress pσ for all investigated steels using a single complex equation σcrR2 = 0.9918 and graphically in R2 = 0.9937, as well as graphically in The excellent accuracy of Equation (26) is documented by the high value of the corresponding coefficient of determination crDRXε on the peak strain pε of all investigated steels is expressed with excellent accuracy by Equations (22)–(26). Equation (24), which expresses the linear dependence of crDRXε on pε of the medium carbon steel C, corresponds very well to the analogous dependence derived in [Linear dependence of the critical strain rived in  for C45 crDRXε on peak strain pε was also derived by Liu et al. [The relationship expressing the dependence of critical strain u et al.  for 316LcrDRXε on peak strain pε for low carbon micro-alloyed steels:The 316LN steel used in  containeThe micro-alloyed steels used in  containeQ decrease with increasing carbon content or with an increasing value of carbon equivalent in the investigated steels, as also documented in C is the carbon content (wt %) and ekvC (−) is the carbon equivalent of the investigated steels. The accuracy of Equations (31) and (32), documented by the relevant coefficients of determination R2 = 0.9062 for Equation (31), or R2 = 0.8813 for Equation (32) is sufficient. Suppose that from graphical dependence of the activation energy on the carbon content or on the size of the carbon equivalent is excluded activation energy of medium carbon steel C, because the deformation temperatures of this steel were chosen by a different methodology than in the present work. Thus, the accuracy of simple prediction of the hot deformation activation energy Q depending on the carbon content will increase significantly Q=−R2 = 0.9894. In the case of Equation (34), the coefficient of determination reaches the value R2 = 0.9999.Excellent accuracy of Equation (33) is documented by the high value of the coefficient of determination Q of similar non-alloyed carbon steels can be found in the literature [Q on the carbon content in non-alloy steels.The values of hot deformation activation energy terature ,64. Thesterature ,64,65,66terature  that theterature ,64, the Parametric equations expressing the dependence of hot deformation activation energy on the chemical composition of steels can also be found in the literature. In work , ElfmarkIn , Medina In , Colas pA comparison of the activation energy values determined by the analysis of the measured data in the ENERGY 4.0 software and calculated using Equations (31) and (35)–(37) are documented in Q with increasing carbon content in the investigated steels. However, the activation energy determined according to Equation (35) is significantly lower in the case of the investigated low carbon steels—see Q—see Equation (36). This equation was determined by analyzing data measured by torsion tests performed for a wide range of non-alloyed and alloyed steels in strain rates 0.54–5.22 s−1 and deformation temperatures 850–1100 °C. The values of activation energy determined according to relation (36) did not differ significantly depending on the carbon content of the non-alloy carbon steels we examined—see Q—see Equation (37). Unfortunately, work [Q value for all the steels examined in this paper. For this purpose, the newly introduced relation (31), which expresses the logarithmic dependence of the activation energy on the carbon content in the investigated steels, seems to be more appropriate.Equations (31) and (35) and the results of the analysis of the measured data in the ENERGY 4.0 software reflect the decrease of the hot deformation activation energy ly, work  does notcrDRXε, can also be used to approximate the flow stress curves of investigated steels. Of course, other models can be used for these purposes, which allow the approximation of the measured flow stress curves in the whole range of strains, such as the model by authors Hensel and Spittel [The model by authors Cingara and McQueen (14) , which w Spittel , etc.pσ, peak strain pε, strain hardening exponent C, or using the relations (4), (5), (6), (14), (20) and the corresponding material constants (see −1. If there is no transition to steady state flow (steady state), the model of Cingara and McQueen (14) can be used to approximate flow stress curves with very good accuracy, not only for the range of strains from 0 up to the pε (or up to stress pσ), but also in the area of the prevailing decrease in flow stress—see In this case, however, it is easier to use the Cingara and McQueen model (14) because the values of all relevant material constants are already determined for all investigated steels. Using the predicted values of peak flow stress ants see , it is, n (−) represents the number of points of the flow stress curves of particular steel included in the calculations . The values of iσ (MPa) then represent the experimentally obtained values of flow stresses and σ(iε) (MPa) are the values of flow stresses obtained by prediction through the model (14). The RMSE value then acquires a specific dimension (MPa). Since the model (14) [pε), the RMSE values were determined only for the predicted flow stress curves in the range of strains from 0 do pε. The excellent accuracy of the predicted flow stress curves (using Equation (14)) of all examined steels in the range of strains from 0 to pε is documented by low RMSE values  ,69:(38)Rdel (14)  was origlues see , which wThe stress–strain curves of selected non-alloy carbon steels were investigated using isothermal continuous uniaxial compression tests performed on a Gleeble 3800-GTC simulator.pσ and the peak strain pε. The critical strains crDRXε required to induce dynamic recrystallization were then determined by analyzing the stress dependence of the work hardening rate on peak stress, using model Cingara and McQueen. In addition, the critical flow stresses crDRXσ were also determined using this model, which corresponds to the critical strains crDRXε of all investigated steels. Based on the achieved results, it was confirmed that for a given size of the Zener–Hollomon parameter, with increasing carbon content  in the investigated steels, their peak flow stress and critical flow stress also increase. In the case of carbon content or carbon equivalent influence, the peak strain and critical strain size made the situation more complicated. However, for all investigated steels, it was confirmed that the difference between the peak strain pε and the critical strain crDRXε increased with decreasing deformation temperature and increasing strain rate.The obtained stress–strain curves were analyzed in the ENERGY 4.0 software to determine the hot deformation activation energy and other material constants suitable for predicting the peak flow stress crDRXε, determined according to Equation (14) and subsequently experimentally verified, equations describing with excellent accuracy the simple linear dependence of crDRXε on the peak strain pε were derived for all investigated steels. In addition, equations comprehensively describing the linear dependence of the critical strain crDRXε on the peak strain pε and the critical flow stresses crDRXσ on the peak flow stress pσ were derived as well.Based on the critical strains Q decreased with increasing carbon content  in the investigated steels, and this dependence was described with good accuracy by a logarithmic equation.The determined hot deformation activation energy pε, was documented by the low determined values of root mean square error.Using the model Cingara and McQueen individual flow stress curves of investigated steels were mathematically described. The excellent accuracy of the predicted flow stress curves of all investigated steels, especially in the range of strains from 0 to


The influence of preparation technology on CaAl–SiO3–LDH structures and properties as well as the flame retardant and smoke suppression characteristics of flame retardant-treated bamboo was discussed. The results revealed that the crystallisation temperature, crystallisation time and crystallisation concentration of CaAl–SiO3–LDHs considerably affected its structure and properties. The optimum technological parameters for preparing CaAl–SiO3–LDHs by using the coprecipitation method are as follows: crystallisation temperature of 100 °C, crystallisation time of 9 h and Ca2+ solution molar concentration of 0.33 mol/L. Compared with nonflame-retardant wood, CaAl–SiO3–LDH flame retardant treatment delayed the peak time of the heat release rate by 20 s and the ignition time by 77.78% and increased the carbon residue rate by 9.54%. This study can provide reference for the research of new flame retardant for bamboo products.Bamboo has been widely used in architecture, decoration and other fields because of its advantages of short growth period, high strength and degradability. However, bamboo, as a combustible material like wood, are easy to burn and cause building fires. However, the existing bamboo water-based flame retardants have some shortcomings, such as strong hygroscopicity and easy loss, which limits the application of bamboo products. In order to improve the flame retardant performance of bamboo, CaAl-SiO China is a major exporter of bamboo products, and more than 60% of popular bamboo flooring, bamboo furniture, planted bamboo and recombinant bamboo exported to Europe and America originates from China. The quality and processing technology of Chinese bamboo products are the global standard . However2PO4)2@HCCP by surface treatment technology to overcome the shortcomings of water-based inorganic flame retardants, layered double hydroxides (LDHs) were developed. LDHs are typical anionic materials with metal hydroxide as the main layer, and anions and some water molecules embedded in the layers to form a layered structure. LDHs provide an easy regulation of physical and chemical properties and rich hydroxyl groups [Currently, phosphorus–nitrogen and nitrogen–phosphorus–boron composites are typically used as bamboo flame retardants. Most of these retardants are inexpensive, water-based inorganic flame retardants that exhibit flame retardant effects, limited pollution and low price . Howeverl groups ,10. Becal groups . Therefo3–LDHs) was prepared using the coprecipitation method and it was used as a flame retardant for bamboo to improve the flame retardant performance of bamboo. The effects of preparation technology on the structure and properties of calcium–aluminium hydrotalcite in root cutting silicate layer and the flame retardant and smoke suppression properties of bamboo treated by CaAl–SiO3–LDHs [The coprecipitation method is a simple and commonly used method to prepare LDHs flame retardant materials ,13,14. LiO3–LDHs  were inv3)2·4H2O), aluminium nitrate (Al(NO3)3·9H2O), sodium silicate (Na2SiO3·9H2O) and sodium hydroxide (NaOH) were obtained from Sinopharm Chemical Reagent Co., Ltd., Shanghai, China. Deionised water was used in the synthesis and washing process. The water was boiled and cooled to room temperature for reuse.Calcium nitrate 2·4H2O and Al(NO3)3·9H2O were weighed and added in the molar ratio of Ca2+: Al3+ = 2 [2SiO3·9H2O was weighed and added to deionised water. Ca2+ solutions with molar concentrations of 0.13, 0.33, 0.53 and 0.73 mol/L were prepared. A high-concentration NaOH solution was prepared for subsequent use. The sodium silicate solution was added to a 500 mL three-neck flask. The flask was placed in a constant temperature oil bath at 25 °C, and the mixture was stirred. The prepared mixed salt solution and NaOH solution were subjected to a constant pressure dropping funnel. The NaOH solution was added until the pH of the solution in the flask reached 10. Next, the mixed salt solution was added to maintain the pH value of the solution at 10. The prepared white slurry was crystallised at 85, 100 and 120 °C and stirred vigorously at a constant speed for 6, 9 and 12 h and subsequently stored at room temperature. Finally, the cake was filtered, washed to neutrality and dried for 24 h at 80 °C. The synthesis mechanism is illustrated in The general formula of hydrotalcite is x < 0.33 . Ca(NO3)Al3+ = 2 ,21,22. A3–LDHs powder was prepared and ultrasonically dispersed in an aqueous solution to obtain 1% suspension, and ultrasonic treatment was performed for 1 h. Then, bamboo chips were placed in the CaAl–SiO3–LDHs suspension. Next, the CaAl–SiO3–LDHs suspension soaked with bamboo chips is placed in a pressurised tank and sealed. Then samples were vacuumised, and when the vacuum degree reaches 0.085 MPa, kept for 20 min. Next, a pressure of 0.6 MPa was applied. When the pressure reached the preset value, the exhaust valve was opened to relieve the pressure to the preset value. The bamboo pieces were soaked for 3 h. Finally, the bamboo chips were removed, the suspension on the surface was removed with deionised water and bamboo chips were dried in an oven to obtain CaAl–SiO3–LDHs flame-retardant bamboo.A CaAl–SiO3–LDHs samples. The X-ray tube was the Cu target, Kα was the radiation source, the tube voltage was 40 Kv, the tube current was 30 mA, the scanning range was 5°–75° (2θ), and the scanning speed was 2 °/min.XRD-6000 manufactured by Shimadzu Corporation of Japan was used to analyse the phase composition and crystallisation of the CaAl–SiO3–LDHs prepared under different conditions, the samples dried in the oven at 80 °C were mixed and ground with potassium bromide in a ratio of 1:100. The synthesised CaAl–SiO3–LDHs samples were analysed using the Prestige-21 FT-IR spectrometer produced by Shimadzu Corporation of Japan, with a scanning range of 400–4000 cm−1.To determine the chemical composition of CaAl–SiO3–LDHs prepared under different conditions were analysed. Cold field emission scanning electron microscope SU8010 produced by Hitachi of Japan was used to analyse CaAl–SiO3–LDHs under an accelerating voltage of 25.0 kV.The morphology and dispersion of CaAl–SiO3–LDHs samples was analysed by using the STA 409C thermogravimetric analyser produced by NETZSCH, Germany. The temperature ranged from 35 to 750 °C, the heating rate was 20 °C/min, the sample mass was 6 to 9 mg and during the reaction process, nitrogen is always introduced.The thermal stability of CaAl–SiO−2, corresponding to temperature of 728 °C [Nonflame-retardant bamboo chips and flame-retardant treaded bamboo chips were bonded with a small amount of white latex into samples with dimensions of 100 mm × 100 mm × 5 mm (length × width × thickness); the combustion performance of the samples was tested with a cone calorimeter , according to the ISO5660-1 procedure. This is to analyse the influences of CaAl–SiO3–LDHs-flame retardants on the flame-retardant effect of bamboo. The heat radiation power was 50 kW/mf 728 °C ,24. Duri3–LDHs prepared by the crystallisation reaction at 85, 100 and 120 °C for 12 h is displayed in The XRD pattern of CaAl–SiO3–LDHs samples synthesised at 85, 100 and 120 °C corresponded to the characteristic crystal planes (003), (006) and (110) of hydrotalcite. The XRD patterns characteristic peak of CaAl–SiO3–LDHs crystallised at 100 °C was the sharpest, and its intensity was the highest, which indicated that the crystal lattice of the sample was the most complete, and the crystallinity was the highest. Therefore, 100 °C is the best crystallisation temperature of CaAl–SiO3–LDHs.Several main characteristic peaks of CaAl–SiO3–LDHs prepared at various crystallisation temperatures are displayed in The FT-IR spectra of CaAl–SiO3–LDHs prepared at various crystallisation temperatures are similar. The absorption peak near 3474 cm−1 was related to the stretching vibration of hydroxyl groups of silicate calcium–aluminium hydrotalcite laminate or hydroxyl groups that physically adsorb water molecules. The absorption peak near 1637 cm−1 was caused by the tensile vibration of hydroxyl groups in crystallised water. The absorption peaks at 1021 and 539 cm−1 were caused by the vibration of silicate ions [−1 was attributed to CO32−,which indicated that CO32− was doped in hydrotalcite, which could be related to the strong adsorption capacity of calcium–aluminium hydrotalcite for CO2 caused by the adsorption of CO2 when the sample was exposed to air after preparation [ate ions , which iparation .3–LDHs prepared at crystallisation temperatures of 85, 100 and 120 °C are displayed in The SEM images of CaAl–SiO3–LDHs samples exhibit obvious layered structure, which shows the characteristic structure of hydrotalcite. The samples prepared at crystallisation temperatures of 85 °C and 100 °C were unevenly distributed and exhibited two morphologies, namely larger lamellar and smaller lamellar. But the samples prepared at crystallisation temperatures of 85 °C, the smaller lamellar structure was present in a larger proportion; among the samples prepared at 100 °C, the smaller layered structure accounts for a large proportion. However, the samples prepared at 120 °C exhibited a uniform morphology distribution, but its lamellar size was typically large.3–LDHs samples prepared at various crystallisation temperatures are displayed in The TG–DTG curves of CaAl–SiO3–LDHs samples crystallised at 85 °C occurred at 165–299 °C, with a weight loss rate of 13%. The weight loss rate and weight loss temperature of the samples crystallised at 85 °C were higher and lower, respectively, than the other two samples, which indicated that the layered structure was unstable and numerous hydroxyl groups were lost. In the third stage of thermal weight loss, the CaAl–SiO3–LDHs samples crystallised at 100 °C exhibited thermal weight loss at 480–564 °C, with a weight loss rate of 5%. Weight loss temperature, the weight loss peak and the thermal stability of the CaAl–SiO3–LDHs samples crystallised at 100 °C were higher, narrower and better than those of the other samples. Therefore, the composition and distribution of interlayer anions in CaAl–SiO3–LDHs samples crystallised at 100 °C were uniform. Thus, 100 °C was the best crystallisation temperature for preparing CaAl–SiO3–LDHs.Thermal weightlessness can be clearly categorised into three weightless stages . The fir3–LDHs prepared by crystallisation reaction at 100 °C for 6, 9 and 12 h are displayed in The XRD patterns of CaAl–SiO3–LDHs sample are similar to the characteristic crystal planes (003), (006) and (110) of hydrotalcite. This phenomenon indicated that the synthesised sample has a typical hydrotalcite lattice structure. The difference in location of each characteristic crystal plane between various samples was not more than 0.2, which indicated that the crystal lattice structure of the synthesised samples was stable when the crystallisation temperature was 100 °C, proving that 100 °C was the best crystallisation temperature for synthesising CaAl–SiO3–LDHs. Among all the samples, the XRD characteristic peaks intensity of the samples with crystallisation time of 9 h was the highest and sharpest, which indicated that the samples with the crystallisation time of 9 h exhibited the highest crystallinity.3–LDHs prepared by crystallisation reaction at 100 °C for 6, 9 and 12 h are displayed in The FT-IR spectra of CaAl–SiO3–LDHs samples were similar , (006) and (110) of hydrotalcite. The XRD patterns of CaAl–SiO3–LDHs synthesised with the molar concentration of Ca2+ of 0.33 mol/L exhibited the sharpest characteristic peak and the highest intensity, which indicated that the sample exhibited the most complete lattice and the highest crystallinity. Therefore, the molar concentration of Ca2+ of 0.33 mol/L is the optimal molar concentration for synthesising CaAl–SiO3–LDHs. Due to the fact that the sample with the concentration of 0.13 does not have the structural characteristics of hydrotalcite, it will not be further explored.3–LDHs prepared with the molar concentrations of 0.33, 0.53 and 0.73 mol/L are displayed in The FT-IR spectra of CaAl–SiO3–LDHs samples are similar. The absorption peak near 3480 cm−1 was related to the stretching vibration of hydroxyl groups in silicate calcium–aluminium hydrotalcite laminate or hydroxyl groups that physically adsorb water molecules, whereas the absorption peak near 1623 cm−1 was caused by the stretching vibration of hydroxyl groups in crystallised water. The vibration peaks of silicate ions appeared at 978 and 539 cm−1, which indicated that silicate ions were successfully embedded in hydrotalcite. The peak of tensile vibration of CO appeared at 1388 cm−1, which indicated that CO32− was doped in hydrotalcite.The positions of FT-IR absorption peaks of three CaAl–SiO3–LDHs prepared with molar concentrations of 0.33, 0.53 and 0.73 mol/L are displayed in The SEM images of CaAl–SiO3–LDHs samples prepared at concentrations of 0.33 and 0.53 mol/L exhibited an obvious layered structure, which is consistent with XRD analysis results. The sample with the concentration of 0.53 mol/L exhibited a uniform appearance, but its size was large. Thus, 0.33 mol/L is the best concentration for preparing CaAl–SiO3–LDHs because the sample with 0.33 mol/L was the most uniform in morphology and size.3–LDHs prepared with molar concentrations of 0.33, 0.53 and 0.73 mol/L are displayed in The TG–DTG curves of CAl–SiO2+ concentration for preparing CaAl–SiO3–LDHs by coprecipitation was 0.33 mol/L.3–LDHs are displayed in The HRR curves of bamboo before and after flame retardant treatment of CaAl–SiO3−LDHs reduced the heat generated by bamboo combustion and inhibited bamboo combustion. When the fire was at its peak and the HRR reached its peak, CaAl–SiO3–LDHs delayed the arrival time and improved the flame retardancy of bamboo.The HRR of materials during combustion can reflect the speed and magnitude of heat released from a fire source . The lar3−LDHs did not easily catch fire, and the flame retardant property of bamboo was considerably improved.The TTI refers to the continuous ignition time required for the flame to burn on the material surface. The longer the ignition time is, the less likely ignition occurs under experimental conditions, and the better the flame retardance of the material is . In the 3−LDHs.3−LDHs improved the carbon residue rate of bamboo, such that CaAl−SiO3−LDHs could promote char formation during combustion.The final remaining mass of nonflame-retardant bamboo chips accounted for 8% of the initial mass, whereas that of bamboo chips embedded with flame retardant accounted for 17.54% of the initial mass. Compared with nonflame-retardant bamboo chips, the residue mass of bamboo after treatment improved by 9.54%, which indicates that the addition of CaAl−SiO(1)3–LDHs considerably influenced its structure and properties. The crystallinity and thermal stability of CaAl–SiO3–LDHs prepared at 100 °C was the highest.The crystallisation temperature of CaAl–SiO(2)3–LDHs considerably influenced its structure and properties. The sample prepared at 9 h exhibited high crystallinity, obvious layered structure and the best thermal stability.The crystallisation time of CaAl–SiO(3)3–LDHs considerably influenced its structure and properties. When the molar concentration of Ca2+ was 0.33 mol/L, the crystallinity of the sample was the highest, the lamellar was the most uniform and the thermal stability is excellent.The molar concentration of solution of CaAl–SiO(4)3–LDHs by the coprecipitation method are as follows: crystallisation temperature 100 °C, crystallisation time 9 h and molar concentration of Ca2+ solution 0.33 mol/L. It is not a cheap raw material, but also has high thermal stability.The optimum technological parameters for preparing CaAl–SiO(5)3–LDHs, the peak times of HRR and ignition were delayed by 20 s and 77.78%, respectively, and the residual carbon rate was increased by 9.54%, the results show that the addition of CaAl–SiO3–LDHs greatly improves the flame retardancy of bamboo.Compared with the nonflame-retardant bamboo chips, after flame retardant treatment with CaAl–SiO