It is well-known that lithium-ion batteries （LIBs） are recognized as one of the most appropriate and promising energy-storage systems，they are also being pursued intensively for transporting applications，including hybrid vehicles （HEVs） and pure electric vehicles （EVs）. However，existing commercial LIBs are the current limitation for the rapid development of both EVs and HEVs due to their limited energy densities and high price cost. Great efforts have been made to develop cathode materials with high voltage，high capacity，low cost，and excellent cycle performance. In this regard，lithium manganate （LiMn₂O₄） with spinel structure has been considered as one of the most promising cathode materials for the next generation high energy density LIBs owing to its high potential （4.1 V），relatively high theoretical capacity （148 mAh·g⁻¹），and abundance of the raw materials. At present，the large-scale commercialization applications of LiMn₂O₄ are plagued by its serious capacity degradation and structural instability during Li⁺ intercalation/de-intercalation process，particularly at higher C rates. Element doping is a widely acceptable modification approach to improve the structural stability of the bulk LiMn₂O₄. Bulk phase doping consists of lithium site，manganese site and oxygen site accor ding to the lattice position occupied by the doped element. Among them，the mainstream of doping studies is manganese site doping，which mainly includes Ni²⁺，Fe²⁺，Zn²⁺，Mg²⁺，Al³⁺，Co³⁺，Ti⁴⁺，etc. The results show that partial substitution of Mn³⁺ can effectively inhibit the structural deformation caused by Jahn-Teller effect，enhance the structural stability of the material，and then improve the cycle performance. Lithium site doping is the partial substitution of the doped element for the Li position in the LiMn₂O₄ structure. Na⁺ is usually used to replace part of Li⁺. Because the electrochemical activity of Na⁺ is worse than that of Li⁺，Na⁺ can remain in the three-dimensional channel of Li⁺ diffusion during the charging and discharging process，and play a role in supporting Li⁺ diffusion channel. Oxygen site doping is the substitution of doping elements such as S，Cl，F，etc. After F replaces O，the bond energy can be enhanced，thereby improving the stability of the crystal structure，and the lattice constant can be increased after Cl and S replace O，facilitating the diffusion of lithium ions. There have been reviews focusing on single element doping for LiMn₂O₄，while the development in the recent years still needs to be summarized，especially the research of double element doping. In this review，the research progress of two-element doped spinel lithium manganate materials in recent years was reviewed，and the future development trend of spinel lithium manganate material was also analyzed.

γ-Ce₂S₃ is a non-toxic and environmentally friendly bright red inorganic pigment with strong color intensity and excellent ultraviolet （UV） resistance，but its oxidation resistance temperature is limited to 350 ℃. In order to enhance the oxidation resistance of γ-Ce₂S₃ red pigment，a coating a layer with high temperature resistant material is usually employed to wrap the pigment. In this paper，a sol-gel method combined with sulfurization technique was used to prepare ZrSiO₄ wrapped γ-Ce₂S₃ red pigment. Ce（NO₃）₃·6H₂O was used as the cerium source，ZrCl₄ as the zirconium source，Si（OC₂H₅）₄（TEOS） as the silicon source，ethylenediaminetetraacetic acid （EDTA） as the complexing agent，HNO₃ as the oxidizing agent，deionized water as the solvent，LiF as the mineralizer，and NH_3·H₂O as the pH regulator. Ce（NO₃）₃·6H₂O and EDTA were dissolved in deionized water in specific proportions to create a uniform solution. ZrCl₄ and TEOS were added to the solution，along with LiF as the mineralizer to lower the formation temperature of ZrSiO₄，promoting its crystallization at 700 ℃. HNO₃ was then added as the oxidizing agent，and the solution's pH was adjusted with NH₃·H₂O to facilitate the complexation of EDTA with Ce³⁺. The prepared solution was stirred to form a uniform precursor solution. The precursor solution was then heated in an oven to evaporate excess water，forming a viscous gel. This gel was subjected to calcination in a muffle furnace，resulting in a porous Ce–Si–Zr–O precursor powder. Using CS₂ as the sulfur source，sulfurization of the precursor powder at 900 ℃ yielded well-developed ZrSiO₄ crystals and transformed CeO₂ into α-Ce₂S₃. Subsequently，the material was heated in a protective argon atmosphere to 1300 ℃，converting it into γ-Ce₂S₃. The influence of different cerium-zirconium ratio of precursor powder and different sulfurization temperatures on the color and morphology of the pigment was investigated. Scanning electron microscopy （SEM，JSM-6700E，JPN） was used to examine the morphology and microstructure of ZrSiO₄ and the coated pigment. X-ray powder diffraction （XRD，D8-Advance，GER） was employed to determine the phase composition of the products，while high-resolution transmission electron microscopy （HRTEM，JEM-2010，JPN） was used to examine the encapsulation of the pigment and the zirconium silicate phase. The thermal analysis （STA449C，GER） was conducted using a thermogravimetric analyzer （TG） and differential scanning calorimetry （DSC）. The color measurements （ACM402，PRC） were carried out using a color analysis instrument，and L*，a* and b* values were used to represent brightness，redness and yellowness，respectively. Fourier-transform infrared spectroscopy （FTIR，Nicolet5700，USA） was employed to characterize the gel transformation and precursor powder at different pre-sintering temperatures. The results showed that the presence of the mineralizer LiF could effectively reduce the formation temperature of ZrSiO₄ crystals，promoting their formation at 700 ℃. However，it also led to the simultaneous generation of Zr₉S₂ and ZrO₂ crystals，which adversely affected the normal color development of the pigment. As the sulfidation temperature increased，Zr₉S₂ and ZrO₂ gradually disappeared by 900 ℃，and they had completely transformed into ZrSiO₄ crystals. Therefore，sulfidation at this temperature was favorable for the color development of the encapsulated pigments. n（Ce）： n（Zr） ratio in the precursor solution significantly influenced the color of the encapsulated pigments. Under the conditions of a sulfidation temperature of 900 ℃，the redness value （a*） of the encapsulated pigments initially increased and then decreased with an increase in the cerium-zirconium ratio. This was because when Zr content was low，the relative content of ZrSiO₄ in the encapsulating phase formed in a unit pigment is insufficient to effectively encapsulate the pigment. On the other hand，an excessive amount of Zr during sulfidation could react with CS₂ to form black Zr₉S_2， leading to darkening of the encapsulated pigments. SEM and TEM images of the encapsulated pigments revealed that they had an inlay encapsulation structure. This structure provided effective protection to pigments embedded in deeper layers，which could greatly enhance the pigments' oxidation resistance. When n（Ce）∶n（Zr） ratio in the precursor solution was 1∶1.0，precursor powder could achieve well-developed ZrSiO₄ crystals during sulfidation at 900 ℃. Additionally，CeO₂ could be vulcanized into α-Ce₂S₃，which could be further transformed into γ-Ce₂S₃ by heating to 1300 ℃ under an argon protective atmosphere. This ultimately resulted in the production with antioxidant temperature of 610.3 ℃ and brightly colored （L*=30.26，a*=37.71，b*=29.24） γ-Ce₂S₃@ZrSiO₄ encapsulated red pigment.

Energy saving and emission reduction have attracted more attention under the background of emission peak and carbon neutrality. As insoluble and environmentally friendly electrodes，the metal oxide electrodes have wide development prospect in the field of power-saved electrolysis，pollutant degradation and electrosynthesis. However，the noble metal oxide electrodes have important problems such as low stability in the strong acidic solution，while the durability of the base metal oxide electrodes is not good in the weak acidic solution. Therefore，this review briefly analyzed the three stages of aging，and the failure mechanism of the noble oxide electrodes. The above stages of aging of the electrodes generally included dissolution of soluble substances，dissolution of active oxides，and detachment of the coating from the titanium substrate surface. It was usually believed that the failure mechanism of the noble oxide electrodes was dissolution of coating，spalling of the coating from the titanium substrate，and the passivation of the titanium substrate. Based on the relationship between the composition，the structure and the properties of the materials，this review overviewed the development of durability for the metal oxide electrodes from three aspects such as the titanium substrate，interlayers and catalytic coating. More specially，in the aspect of the titanium substrate，the cementation technology was simple but relatively limited for the catalytic coatings prepared by thermal decomposition. The composition of the electrode interlayers went through refractory metals，semiconductor metal oxides，and metal matrix composites. The preparation methods of the above electrode interlayers included electroplating，anodization，micro-arc oxidation，thermal spray，cold spray and thermal oxidative decomposition. The optimal method usually depended on the electrode interlayer composition and production cost. In general，the catalytic coating of the electrodes might be modified by elements doping，particle addition，and regulation of the deposition behavior of metallic oxides. In addition，it analyzed the adaptability of the electrodes under harsh condition such as pulse electrolysis，periodic reversing electrolysis，electrolysis in the low temperature solution，and liquid flow enhanced electrolysis. The metal oxide electrodes modified through coating composition and structure optimization，had a good ability to adapt to extreme environment. Finally，this review suggested the inert component choice and structure regulation should be studied for the noble metal oxide electrodes，while the doping modification and coating system design should be worth exploring for the base metal oxide electrode.

With the increase of energy consumption and the depletion of resources，developing and finding alternative energy sources is the key to alleviate these problems. Proton exchange membrane fuel cells （PEMFC） is widely used in automotive field with its advantages of green and high resource utilization，which is one of the ultimate solutions for clean energy vehicles and has attracted much attention in recent years. PEMFC is an ideal power source for new energy vehicles and has attracted much attention in recent years. Usually，PEMFC components are mainly composed of membrane electrode assembly，bipolar plate and seals，of which the bipolar plate is one of the key multifunctional components in PEMFC，which accounts for about 80% of the total mass of the stack and 45% of the cost. Therefore，the preparation of low-cost，small size，high-performance bipolar plate is decisive for PEMFC power density，service life，etc. Currently，the bipolar plate can be used as an electrode for PEMFC，which can be used for PEMFC. At present，the materials that can be used as bipolar plates are graphite bipolar plates，metal bipolar plates and composite bipolar plates，and the most promising is the metal bipolar plate. Among them，titanium bipolar plate has high specific strength，corrosion resistance and electrical conductivity，and has great application prospects in PEMFC bipolar plate. Due to the presence of a certain amount of F⁻ and H⁺ in PEMFC environment，it is difficult for the unmodified titanium bipolar plates to meet the requirements of fuel cells，and the poor mechanical properties of titanium lead to the difficulty of guaranteeing the bipolar plates in terms of the forming quality and the assembly accuracy. Therefore，it is urgent to modify its surface and optimize its forming process. This review firstly introduced the characteristics of the three commonly used bipolar plate materials，and then mainly described the conductive corrosion resistance of titanium metal in PEMFC environment，and it was found that TC4 and TA2 had poorer corrosion resistance and higher interfacial contact resistance in the acidic and F⁻containing environment. In the surface modification of titanium bipolar plates，the advantages and disadvantages as well as the application prospects of surface modification technologies such as physical vapor deposition （PVD），chemical vapor deposition （CVD），electroplating and chemical plating，gas nitriding and plasma nitriding were reviewed. These modification processes enabled the preparation of carbon coatings，transition metal carbide coatings，transition metal nitride coatings，noble metal coatings，conductive polymers and other coatings. Among them，PVD technology had the advantages of non-polluting environment，low film-forming temperature，high film-forming efficiency，etc.，which had a greater potential for application in the surface modification of metal bipolar plates compared with several other modification methods. However，monolayer coatings prepared by conventional PVD technology still suffered from many defect problems such as holes and cracks. Therefore，many researchers had chosen to prepare multilayer composite coatings on the surface of metal bipolar plates so that to reduce the pathway of corrosive fluid intrusion，or to carry out elemental doping in the coatings so as to inhibit the generation of transition metal nitrides，which were better solutions to such problems. In addition，with the continuous development and innovation of film-forming technology，there were many new PVD technologies，these technologies could give the ions higher energy and ion ratio，so that the preparation of the monolayer coating was dense and uniform，with good bonding，which was currently one of the ideal coating preparation processes. In the forming of titanium bipolar plates，the differences in mechanical properties between titanium and stainless steel were compared. Due to the poor mechanical properties of titanium metal，poor forming limit and large rebound deformation，the traditional forming method was no longer suitable for titanium bipolar plates，so there was an urgent need to optimize and innovate the forming process of titanium bipolar plates. Therefore，this paper reviewed the characteristics of electromagnetic forming，three-dimensional （3D） printing forming，multi-stage stamping forming and rubber forming and their development prospects from the perspectives of forming limit and forming performance of titanium bipolar plates. It was found that all four optimized forming methods were able to form ultra-thin titanium sheets without fracture，deformation or rebound of the formed sheets，which could make the formed bipolar plate can meet the demand of automobile light weight while satisfying the performance，which was of great significance for the future development of titanium bipolar plate forming process. At present，the main problems faced by titanium bipolar plates were still the surface modification process and forming process. Surface modification should start from the cost of preparation process，cycle time and coating structure of titanium bipolar plate surface coating，using low-cost coating materials （such as amorphous carbon，metal carbon nitride） ，through the design of multi-layer structure，reduce the total thickness of the coating，to achieve the effect of improving the coating performance and shorten the preparation cycle. As for the forming process，it mainly focused on the forming of titanium bipolar plates. As the precision and complexity of the bipolar plate flow field increased，the demand for titanium forming capabilities continued to rise，so it was necessary to carry out research from the alloy design of titanium alloys and the forming process （method，temperature，lubricant） ，etc.，to improve the forming ability and dimensional accuracy of titanium metal bipolar plates.

WE43A rare earth magnesium alloy possesses advantages in terms of high specific strength，specific stiffness，and high temperature resistance，making it highly beneficial for lightweighting aerospace applications. It finds extensive use in aviation，aerospace，and related fields，such as helicopter frames，aircraft wings，and compressor casings. According to statistics，the majority of magnesium alloy components （90%） are produced through casting，which inevitably leads to the occurrence of porosity defects within the magnesium alloy. Even a small amount of porosity defects can cause a significant decrease in mechanical performance. This is due to the reduction in the effective load-bearing area of the castings caused by the presence of porosity defects，resulting in stress concentration and the initiation and propagation of fatigue cracks under cyclic loading conditions. As a result，the fatigue performance of critical load-bearing castings is compromised，leading to premature failure and severe consequences. In general，the generation of defects can be minimized by implementing techniques such as centrifugal casting and controlling the solidification temperature. Centrifugal casting promotes grain growth，thereby reducing defects associated with grain boundaries and producing a more uniform and dense material structure. Additionally，controlling the solidification temperature allows for the adjustment of the solidification rate，preventing defects caused by excessively high or low local temperatures and improving the overall quality of the material. However，there has been limited systematic investigation on the mechanical properties of materials with defects. Therefore，it is necessary to analyze and study the mechanical properties and changes of WE43A rare earth magnesium alloy containing porosity defects. In this study，WE43A rare earth magnesium alloy was selected as the research subject，and the samples were graded and screened to obtain specimens that meet the defect level requirements for subsequent mechanical property testing. This was achieved by increasing the probability of defect occurrence through process adjustments and directly selecting samples from a large number of production products to obtain specimens that conform to the defect level requirements. X-ray non-destructive testing was an efficient and accurate quality inspection method that played a crucial role in magnesium alloy castings. This technique provided internal structural imaging information of the samples，enabling comprehensive assessment of internal quality. According to the "Specification for Magnesium Alloy Castings"，five defect levels （I to V） were defined to describe the severity of porosity defects. Sponge-like and dispersed forms were two typical types of porosity defects. Four-point bending fatigue tests were performed on specimens with dimensions of approximately （7.4±0.1） mm in thickness，（21.0±0.1） mm in width，and （82.2±0.1） mm in length. The fatigue specimens were plate-shaped，supported by rollers，with a roller diameter of 10 mm，and a pressure roller was used to apply an average stress of 100 MPa at a loading frequency of 20 Hz. Tensile tests were conducted on rod-shaped specimens with a diameter of approximately （5.00±0.02） mm and a length of （110±0.1） mm using an MTS electronic universal testing machine. The tensile tests were carried out in accordance with HB5143，with a strain rate of 0.00025 s⁻¹. Metallographic analysis of the alloy was performed using an XJP-3Ad metallographic microscope （OM）. Scanning electron microscopes （SEM），including S-3400N-type and SU1510-type，were utilized to observe the fracture morphology. The influence of different types and degrees of defects on the mechanical properties of rare earth magnesium alloys was systematically investigated，and the formation mechanism and control strategies of defects were discussed. This provided a theoretical basis for improving the casting quality and efficiency of rare earth magnesium alloys. The study of internal defects in alloys and further understanding of the destructive effects of defects had both theoretical and practical significance. The results indicated that the porosity defects mainly existed as irregular strip-like and elliptical voids at grain boundaries. With increasing defect levels，the area of porosity defects within the specimens continuously increased from 0.8% to 30%. The tensile strength，yield strength，and elongation of WE43A magnesium alloy exhibited a decreasing trend with increasing defect levels，with the decrease in elongation being slightly smaller than that in tensile strength. The average tensile strength of specimens with defect Level Ⅰ was approximately 260 MPa，which decreased to 80 MPa when the defect level reached Ⅴ. The yield strength decreased from 180 to 70 MPa. The reductions in tensile strength and yield strength were 70.71% and 62.22%，respectively，while the decrease in elongation was 41.56%. The average fatigue logarithmic life （lgN） of specimens with defect Level Ⅴ decreased by 28.17% compared to that of defect-free specimens. The average fatigue limit of specimens with defect Level Ⅴ decreased by 38.09% compared to that of defect-free specimens. The fatigue fracture exhibited cleavage feature；a small amount of porosity （defect Level Ⅰ） had a minor effect on the mechanical properties of the material. When the porosity content was low，there was little difference in tensile strength，yield strength，elastic modulus，and defect-free magnesium alloy. However，a noticeable decrease in mechanical performance occurred when the defect level reached Ⅱ. The excessive porosity defects significantly reduced the material's ductility.

Zr-2.5Nb pressure tube is the core structural components of the heavy water reactor nuclear reactor cores，which needs to be replaced every 20 years to ensure the safety，and it's time for Qinshan heavy water reactor power plant to change. Meanwhile，the manufacturing process of Zr-2.5Nb pressure pipe is complicated and the control parameters are numerous，e.g.，hot extrusion，cold draw etc.，it's difficult to control the final microstructure and the corresponding properties. As a components of nuclear reactor cores，Zr-2.5Nb pressure tube are subjected to harsh service environments and stringent performance requirements，e.g.，high strength，plastic toughness，good high temperature performance，and corrosion resistance to irradiation，etc. Thermal deformation involving hot forging and hot extrusion，is not only a necessary process for Zr-2.5Nb pressure tube to obtain the designed shape，but also an effective method to change the microstructure （phase composition，grain size/topography/orientation，substructure，etc.） of materials. Hot deformation plays a crucial role in the engineering applications and scientific research，due to the controlling microstructure and expected performance of the alloy by the process of work-hardening and dynamic softening with the core transformations of phase change，recrystallization and grain growth. Therefore，analyze the various types of phenomena that may occur during hot deformation and elucidating the corresponding microscopic mechanism of the phenomena can lay a theoretical foundation for the design of the process and the regulation of the organization in the actual production. The thermal processing diagrams take into account the effect of thermal deformation parameters，i.e.，deformation temperature T，strain rate $\dot{\epsilon }$ and strain ε，on the thermal deformation behavior of the material and provide a preferred processing range for engineering applications. Combine with the microstructure of the typical peak energy consumption region，the thermal deformation behavior of the alloy and its mechanism are analyzed，finally achieve the purpose of controlling the microstructure，avoiding defects and optimizing the process parameters. In this paper，the thermal deformation behavior of Zr-2.5Nb alloy was analyzed by using Gleeble 3500 to conduct thermal compression tests under different conditions （i.e.，750~960 ℃，0.01~10 s⁻¹） ，constructing the thermal processing map of the alloy，and characterizing the typical microstructure by electron scanning electron microscope （SEM） and electron backscatter diffraction （EBSD） . The results showed that the alloy mainly underwent α-phase basal slip at lower temperatures （T<810 ℃） ，which led to strong grain selective orientation due to grain rotation and a large number of defects resulting in deformation instability. In addition，the alloy existed in two preferred processing zones with high energy consumption values and stable deformation： the medium-temperature preferential processing zone （810 ℃<T<840 ℃，$\dot{\epsilon }$<0.1 s⁻¹，ε>0.5） ，where deformation and dynamic recrystallization occurred simultaneously of α-phase and β-phase in the alloy，the room-temperature organization consists of polygonal/long-stripe α-phase and martensite α'；high temperature preferred processing zone （850 ℃<T<900 ℃，ε>0.5）， the work hardening and dynamic softening balance in this zone，due to the simultaneous occurrence of deformation and dynamic recrystallization of β-phase，resulting in the complete formation of α' martensite in the room temperature. In summary，Zr-2.5Nb alloy was unstable in deformation at low temperature （T<810 ℃），and at high temperature （T>850 ℃） the organization was β-phase with large grain size resulting in α' martensite after quenching，in order to achieve the purpose of regulating the microstructure of the alloy，the temperature and the strain rate could be controlled in the medium-temperature preferred machining zone （810 ℃<T<840 ℃，$\dot{\epsilon }$<0.1 s⁻¹，ε>0.5） .

In order to investigate the effects of oxidation conditions on the leaching rate of niobium tantalum ore，the cyclic voltammetry，Tafel polarization curve，impedance and open-circuit potential methods were used to study the effects of different oxidation potentials，different HF concentrations and different temperatures on the leaching rate of niobium tantalum ore. The cyclic voltammetry was used to investigate the oxidation reaction in the process of acid leaching of niobium tantalum ore. The cyclic voltammetry curve of niobium tantalum electrode had four oxidation peaks A₁，A₂，A₃ and A₄ of low valence Fe and Mn，and the oxidation peak potential values were -0.38，0.1，0.31 and 0.6 V，respectively. Among them，A₁ and A₂ were the oxidation peaks of MnO and Fe₃O₄ phases in niobium tantalum ore，corresponding to the oxidation of Mn²⁺ to Mn³⁺ and Fe²⁺ to Fe³⁺，respectively. A₃ and A₄ were the oxidation of （Fe，Mn）（Nb，Ta）₂O₆ mineral phase in niobium tantalum ore corresponding to the oxidation of Mn²⁺ to Mn³⁺ and Fe²⁺ to Fe³⁺ in （Fe，Mn）（Nb，Ta）₂O₆ mineral phase，respectively. The current density （J_coor） of Tafel polarization curve increased with the increase of HF concentration in the electrolyte，and reached the maximum value of 849.68 µA·cm⁻² at 10 mol·L⁻¹ HF. The anode slope decreased with the increase of HF concentration and then tended to be stable，indicating that the oxidation leaching rate of total mineral phase at the ore/liquid interface on the surface of niobium tantalum electrode increased with the increase of HF concentration and then tended to be stable after reaching the maximum value at 10 mol·L⁻¹ HF. At -0.38 and 0.1 V oxidation potential，AC impedance test of HF concentration on MnO and Fe₃O₄ mineral phases showed the charge transfer density （J_Rct） increased with increasing HF concentration. Mn charge transfer density J_Rct increased from 3.86 µA·cm⁻² of 4 mol·L⁻¹ HF to 68 µA·cm⁻² of 10 mol·L⁻¹ HF. Fe charge transfer density J_Rct increased from 1.11 µA·cm⁻² of 4 mol·L⁻¹ HF to 87.32 µA·cm⁻² of 10 mol·L⁻¹ HF. The leaching rates of Mn and Fe also accelerated with the increase of HF concentration in the electrolyte. Mn leaching rate reached the maximum of 29.4 mg·L⁻¹·h⁻¹ at 10 mol·L⁻¹ HF and Fe leaching rate reached the maximum of 5.5 mg·L⁻¹·h⁻¹ at 10 mol·L⁻¹ HF. The experimental results showed that the leaching rates of MnO and Fe₃O₄ mineral phase increased with the increase of HF concentration，and the leaching rates of Mn and Fe ions reached the maximum in 10 mol·L⁻¹ HF electrolyte. At 0.31 and 0.6 V oxidation potential，AC impedance test of HF concentration on （Fe，Mn）（Nb，Ta）₂O₆ mineral phase showed that Fe and Mn charge transfer density J_Rct increased with the increase of HF concentration in the electrolyte. Fe charge transfer density J_Rct increased from 1.94 µA·cm⁻² of 4 mol·L⁻¹ HF to 143.7 µA·cm⁻² of 11 mol·L⁻¹ HF. Mn charge transfer density J_Rct increased from 13.06 µA·cm⁻² of 4 mol·L⁻¹ HF to 168.4 µA·cm⁻² of 11 mol·L⁻¹ HF. The leaching rates of Fe and Mn element from niobium tantalum was also accelerated with the increase of HF concentration in the electrolyte. The leaching rate of Fe reached the maximum of 2.56 mg·L⁻¹·h⁻¹ at 11 mol·L⁻¹ HF and the leaching rate of Mn reached the maximum of 2.88 mg·L⁻¹·h⁻¹ at 11 mol·L⁻¹ HF. The experimental results showed that the leaching rate of （Fe，Mn）（Nb，Ta）₂O₆ mineral phase increases with the increase of HF concentration，and the leaching rate of Nb ion reached the maximum at 11 mol·L⁻¹ HF. Tafel results of niobium tantalum ore at different reaction temperatures showed that temperature had a significant effect on the oxidation leaching of niobium tantalum ore. With the increase of temperature，the oxidation leaching rate of niobium tantalum ore was accelerated，and the corrosion current density gradually increased. The corrosion current density tended to flatten out when it reached the maximum value of 959.20 µA·cm⁻² at 80 ℃. The dominant activation energy of Nb ion in the oxidation leaching stage was 3.018 kJ·mol⁻¹ based on Arrhenius formula. The results of AC impedance tests of （Fe，Mn）（Nb，Ta）₂O₆ mineral phase at different reaction temperatures and the determination of Nb ion leaching rate by inductively coupled plasma mass spectrometer （ICP） showed that with the increase of HF concentration，the oxidation current density also increased gradually，reaching the maximum 178.4 µA·cm⁻² and 192.5 µA·cm⁻² at 80 ℃. As determined by ICP，the oxidation leaching rate of Nb accelerated with the increase of temperature，reaching a maximum of 5.32 mg·L⁻¹·h⁻¹ at 80 ℃. The oxidation reaction rate of Nb ion was basically unchanged as the temperature continues to rise，which might be caused by the change of the double layer structure at the niobium tantalum/liquid interface with the increase of temperature. The results of impedance AC experiments at different temperatures showed that the current density increases with the increase of temperature，reaching the maximum values 192.5 µA·cm⁻² at 80 ℃.

Group 5B metals （Nb，V and Ta） and related alloys are considered as promising hydrogen permeation membranes to replace commercial Pd and its alloys. Vanadium-based alloys，as a new substitute for Pd membranes，not only have lower cost but also higher hydrogen permeability. Despite all this，huge challenges remain-especially the poor mechanical stability and high susceptibility to hydrogen embrittlement （HE） which is closely related to high absorbed hydrogen concentration of vanadium. Among group 5B alloys，hydrogenation （thermal-hydrogen treatment technology） has obvious advantages in improving the properties of V-based alloys.However，no relevant research has been carried out on V-TiFe alloys so far. Based on this，the hydrogenation behavior of （Ti₆₅Fe₃₅）_100-xV_x （x=0，2.5，5，10，15，25），x is the percentage of atoms of V） alloys with low cost and high hydrogen permeability were explored. In the solid state，two operating processes of static hydrogenation and dynamic hydrogenation were used，besides，the effects of hydrogen flow，hydrogenation time and hydrogenation temperature （350 and 450 ℃） on the hydrogenation behavior of the above alloys with different V content were systematically studied by the control variable method. The solidification path，microstructure evolution and mechanical properties of the alloy were analyzed by thermodynamic calculation software （Thermo-Calc），scanning electron microscope （SEM），energy spectrometer （EDS），X-ray diffraction meter （XRD），and Vickers hardness tester. First，the composition-temperature curve obtained from the thermodynamic solidification data of Thermo-Calc using the origin software found that： when 0<x<8，the primary phase was TiFe；when x=8，the alloy appeared fully eutectic；and when 8<x<25，the primary phase was body-centered cubic （bcc）-V.With the decrease of solidification temperature，binary and ternary eutectic reactions occurred successively until the end of solidification. Subsequently，the phase analysis of （Ti₆₅Fe₃₅）_100-xV_x cast alloys by XRD，SEM and EDS showed that： when 0<x<10，the alloy was composed of the primary TiFe phase and {bcc-（V，Ti）+TiFe} phase；when x=10，the alloy consisted of a divorced eutectic {β-Ti+TiFe}；when x>10，the alloy consistd of a primary bcc-（V，Ti） solid solution phase and a small amount of TiFe phase. In addition，exploring the changes of hydrogenation amount of （Ti₆₅Fe₃₅）_100-xV_x alloys caused by the difference in hydrogen flow，hydrogenation time and hydrogenation temperature，it was found that： when V content was constant，the hydrogenation effect of alloys was proportional to the variable parameters under any single variable parameter such as hydrogen flow，hydrogenation time and hydrogenation temperature，in other words，when the hydrogenation temperature and hydrogen flow kept a constant，the hydrogenation amounts of the alloy grew with the increase of hydrogenation time. After the hydrogenation time reached a certain level，the hydrogenation amount reached the maximum and tends to be stable，the alloy composition also had a decisive influence on the hydrogenation amount of the alloy besides，that was to say，it would cause an increase in hydrogenation with the increase of V content. Finally，the mechanical properties of all alloys of the system were studied，and the hardness of the cast microstructure of （Ti₆₅Fe₃₅）_100-xV_x cast alloys and the hardness after hydrogenation at 350 and 450 ℃ were measured using Vickers hardness tester. The hardness of the alloy of the hydrogenated V-TiFe system was significantly enhanced compared with the alloy sample before hydrogenation，that was，Vickers hardness of （Ti₆₅Fe₃₅）₈₅V₁₅ alloy after hydrogenation at 450 ℃ was HV_0.3 720.4，which was equivalent to 1.12 times before hydrogenation，and the hardness increased with the increase of hydrogenation temperature，which could be attributed to the solid solution strengthening of H atom inside the alloy and the precipitation of hydrides.

Nonferrous metals，which are the basic materials for the development of China's national economy，are widely used in navigation，aviation，electronic communication，machinery manufacturing industry and many other fields due to their unique physical，chemical and mechanical properties that not provided by steel materials. The South China Sea is a significant maritime transportation hub and military hub of China，possessing a powerful political position. Thus，a large number of nonferrous metal materials are served in the South China Sea and has played a vital role in the construction of ships，deep submergence equipment，offshore drilling platforms，offshore helicopters and other fields. In the 21st century，intensifying the engineering construction in marine has become a significant part of the strategy of national economy development，so the South China Sea has become one of the most important service environments for many marine equipment. However，the corrosion and failure of materials in the marine environment seriously threaten the service safety of engineering equipment. For various structural materials，marine corrosion will not only cause damage and loss of functions of various infrastructures，equipment and structures，shorten the service life of materials and structures，cause huge waste of resources，materials and energy，but also lead to sudden disasters and accidents，and even cause casualties. The South China Sea has a typical harsh marine environment. The environmental characteristics of high temperature，high humidity，high salt，and high sunlight radiation all year round make the non-ferrous metals serving here extremely corrosive，which often causes losses that cannot be ignored. Nonferrous materials have various types with different electrode potential，crystal structure and surface properties，so they have different corrosion characteristics and development laws. For example，aluminum and its alloys are relatively corrosion resistant due to the existence of dense oxide film on its surface，and the main corrosion form is pitting corrosion. The magnesium and its alloy have poor corrosion resistance due to the low electrode potential of magnesium and their corrosion is a dynamic process from pitting corrosion to uniform corrosion；Titanium alloy has strong corrosion resistance，and its performance is mainly affected by biological fouling. Nickel based alloys have high corrosion potential and strong corrosion resistance in the South China Sea environment，and their main corrosion in seawater is pitting corrosion.On the other hand，the corrosion behavior of the same metal in different marine environmental zones is also different.The marine environment can be divided into five zones，namely，marine atmospheric zone，spray splash zone，tidal range zone，seawater immersion zone and sea mud zone. Among them，the marine atmospheric area is the main service environment for non-ferrous metals，and the non-ferrous metal corrosion in this environment is under the collaborative control of land atmosphere，sea salt particles，solar radiation and other factors. The metal materials in the spray splash area have the most serious corrosion due to the alternating dry and wet environment.The corrosion of metals in tidal range area is relatively light due to the effect of seawater tide. The metal corrosion in seawater immersion area is mainly controlled by dissolved oxygen concentration，salinity，temperature and other factors in seawater，while the metal corrosion in sea mud area is mainly caused by microorganisms. According to the research status of corrosion behavior of various nonferrous metals in the South China Sea environment，this review analyzed the corrosion characteristics，rules and hidden dangers of aluminum，titanium，copper，magnesium，nickel，zinc and their alloys in the specific environmental zone of the South China Sea，and summarized the corrosion protection measures of aluminum，titanium，copper，magnesium applicable to the special environmental conditions of the South China Sea. At the same time，some deficiencies in the current research on the corrosion of nonferrous metals in the South China Sea environment were pointed out： 1） For the explanation of metal corrosion mechanism，most of them only focused on the external environmental factors，and there was a lack of in-depth and extensive research on the reasons that metal body was easy to induce certain corrosion forms，such as crystal structure，alloy composition and so on. 2） The corrosion research mainly focused on the marine atmospheric area，and there was little research on the corrosion behavior in the deep-sea environment，so a complete system could not be formed. In the future，the application of nonferrous metals and their alloys was expected to further develop in the deep-sea field，and the research on their deep-sea corrosion mechanism had far-reaching significance for the development of deep-sea engineering. 3） There was a lack of research on anti-corrosion strategies for tropical marine environment. Most non-ferrous metal marine anti-corrosion technologies were tested in simulated seawater，and their anti-corrosion life in the extreme marine environment of the South China Sea might be shortened or even invalid.Therefore，it was necessary to strengthen the research on nonferrous metal anti-corrosion strategies for the tropical marine climate of the South China Sea，and conducted inspection in the real sea environment to obtain more accurate results.

Cyanide wastewater produced in cyanide gold extraction process is complicated in composition and difficult to treat. Direct discharge without treatment will cause great harm to the environment. The existing cyanide wastewater treatment methods generally have the problems of high energy consumption and immature development. Therefore，the search for an efficient，environmentally friend and energy-saving treatment method has aroused widespread concern of researchers. Persulfates have the advantages of high stability，easy operation and low price. Its activation can produce SO₄^·- with a long half-life （40 µs），a wide range of pH application，and strong oxidation，which has attracted extensive attention of researchers. In this paper，sodium persulfate was used as oxidant to treat cyanide wastewater，and the effect of sodium persulfate addition on the removal rate of cyanide and heavy metal ions and the catalytic oxidation reaction mechanism of the process were mainly studied. By means of flame atomic absorption spectrometer （AAS），X-ray diffraction （XRD） and quenching experiment，the changes of metal ions concentration，the types of main precipitation，the types of free radicals and their contributions during the oxidation process were systematically analyzed. In addition，the change of pH and the transformation of S₂O₈²⁻ with the increase of sodium persulfate addition were also discussed. The results showed that the initial pH value of the water sample was 12，the concentration of total cyanide （CN_T） was 2862.2 mg·L⁻¹，the free cyanide （CN⁻） was 1691.3 mg·L⁻¹，Cu⁺ was 789.27 mg·L⁻¹，Fe²⁺ was 42.197 mg·L⁻¹，Zn²⁺ was 131.047 mg·L⁻¹. With the increase of sodium persulfate addition，the removal of CN_T and Cu⁺ firstly increased and then decreased，and the removal rate of CN⁻ and Fe²⁺ soon reached the maximum and remained unchanged. This was because the strong alkalinity of cyanide wastewater and SO₄^·− produced by persulfate activated by Cu⁺ released from the breaking of metal cyanide complex ions strengthen the oxidation of the system and promote the degradation of cyanide. When the persulfate concentration was 0.084 mol·L⁻¹，the maximum removal rates of CN_T，CN⁻，Fe²⁺ and Cu⁺ were 89.2%，100%，97.4% and 92.7%，respectively. When the persulfate concentration increased to 0.105~0.125 mol·L⁻¹，the concentration of sodium persulfate was so high that the excessive S₂O₈²⁻ in the system became the scrubber of SO₄^·−. At the same time，excessive free radicals also appeared self-quenching phenomenon，the oxidation capacity was weakened，and the removal rates of CN_T and Cu⁺ began to decline. In the oxidation process，the metal cyanide complex ions in the system were removed in the form of precipitation. After the addition of sodium persulfate，the weak complex Cu（CN）₃²⁻and Zn（CN）₄²⁻ occurred the detoxification reaction of Cu⁺ and Zn²⁺，CN⁻ and the strong oxidizing S₂O₈²⁻ reaction was quickly consumed. Fe（CN）₆⁴⁻ was removed from the solution by coprecipitation reaction with the dissociated Cu⁺ and Zn²⁺. At the persulfate concentration of 0.042 mol·L⁻¹，CuCN and Zn₂Fe（CN）₆ were the main precipitates. With the increase of the persulfate concentration，CuCN dissolved gradually. Finally，the system formed Zn₂Fe（CN）₆，Cu₂Fe（CN）₆ and CuSCN precipitates，and cyanide ions were completely oxidized to CO₂ and N₂. As the reaction proceeded，the pH of the system gradually decreased，because S₂O₈²⁻ would hydrolyze and consume H₂O or OH⁻ in the system，resulting in the increase of H⁺ concentration in the solution. When the addition of persulfate was different，the concentration of S₂O₈²⁻ decreased gradually with the passage of reaction time. S₂O₈²⁻ was almost completely consumed in wastewater for about 5 min，this was because the persulfate oxidation process was very fast. Both direct oxidation and free radical oxidation consumed S₂O₈²⁻，the added S₂O₈²⁻ was completely converted into SO₄²⁻ and stored in the solution after the reaction. Methanol （MeOH） and tert-butanol （TBA） were added into the persulfate system as quenchers of active species to investigate the species of free radicals in the system. When tert-butanol was added into the system to quench the effect of ·OH，CN_T removal rate decreased from 89.2% to 87%. When methanol was added，the removal rate of CN_T decreased to 77.3% because it could simultaneously quench SO₄^·− and ·OH，which proved the existence of SO₄^·− and ·OH radicals in the system. The removal process of cyanide and metal ions in cyanidation wastewater was the result of the combined action of direct persulfate oxidation and free radical oxidation of SO₄^·− and ·OH produced by activation. The contribution of direct persulfate oxidation was 86.6%，while that of free radical oxidation was only 13.4%.