Corrosion-resistant and high-temperature-resistant copper alloy pipe and preparation method thereof

Abstract

The application discloses a kind of corrosion-resistant high-temperature copper alloy pipe and preparation method thereof, by P, X, copper and inevitable impurities, X is selected from at least one of Fe, Mn, Ti, Cr, Zr, Co, with mass percentage content, 0.12%≤P≤0.36%, 0.0012%≤Fe≤0.50%, 0.0015%≤Mn≤0.45%, 0.0012%≤Cr≤0.40%, 0.0015%≤Ti≤0.30%, 0.0013%≤Zr≤0.27%, 0.0015%≤Co≤0.15%, X total content meets 0.0012%≤X≤0.50%, inevitable impurities total amount≤0.01%, copper is remainder.The application is by synergic control X and the addition ratio of P, optimizes the microstructure of alloy, and the copper alloy pipe material with excellent thermal stability, corrosion resistance and processability is prepared, and the pipe material has wide application prospect in heat exchange, new energy, high-power computer and other fields.

Description

Technical Field

[0001] This application relates to the field of non-ferrous metal materials and their processing technology, specifically to a corrosion-resistant and high-temperature resistant copper alloy tube and its preparation method. Background Technology

[0002] Phosphorus-deoxidized copper tubing (TP2, phosphorus content 0.015~0.04%) possesses excellent thermal conductivity, ductility, and weldability, making it one of the most widely used copper processing materials globally. However, with increasingly complex service environments, TP2 copper tubing has gradually revealed two fatal flaws, severely restricting its application in high-end fields.

[0003] I. Corrosion failure

[0004] In environments containing chloride ions (such as tap water and seawater) and in humid environments containing organic acid vapors such as formic acid and acetic acid, TP2 copper tubes are prone to anthole corrosion, which eventually leads to perforation and leakage. Anthole corrosion is the most common failure mode of heat exchange tubes. According to statistics, about 70% of copper tube leakage accidents in heat exchange systems are caused by anthole corrosion.

[0005] II. Abnormal grain growth after brazing

[0006] During the assembly of heat exchange tubes, brazing is required. The welding temperature is typically as high as 800-1000℃. After welding, the TP2 copper tubes undergo severe abnormal grain growth, with grain sizes reaching 500-700 μm, 50-100 times the original grain size. This abnormal grain growth leads to a significant decrease in the strength, plasticity, and toughness of the weld joint, making it prone to cracking and failure during subsequent assembly and long-term operation under alternating hot and cold temperatures. Statistics show that approximately 20% of copper tube leaks in heat exchange systems occur at the weld joint.

[0007] To address the aforementioned issues, existing technologies primarily improve the corrosion resistance and high-temperature grain size stability of copper alloys by adding alloying elements.

[0008] CN103866157B discloses a corrosion-resistant and high-temperature-resistant copper alloy tube, which improves the recrystallization temperature and corrosion resistance of the alloy by adding 0.2~1.0% Sn and 0.005~0.05% La / Ce to TP2. However, the addition of Sn significantly reduces the thermal conductivity of the copper alloy; rare earth elements are expensive, highly reactive, difficult to control in smelting, and easily form oxide inclusions, reducing the alloy's machinability.

[0009] CN117242197A discloses a corrosion-resistant copper alloy. The corrosion resistance of the alloy is improved by adding metal elements (Li, K, Ca, Na, Mg) with a standard electrode potential lower than that of Al. However, these elements have high chemical activity and are easily oxidized and volatilized during the high-temperature brazing process, which damages the alloy's processing performance, weakens the bonding strength between the matrix and the brazing filler metal, and may accelerate high-temperature grain coarsening due to grain boundary segregation.

[0010] Given the shortcomings of existing technologies, there is an urgent need to develop a copper alloy tube that combines excellent high-temperature brazing compatibility with resistance to ant hole corrosion in order to meet the needs of high-end fields such as heat exchange, new energy, and high-power computers. Summary of the Invention

[0011] The purpose of this invention is to overcome the shortcomings of existing technologies and provide a corrosion-resistant and high-temperature resistant copper alloy tube and its preparation method. The copper alloy tube of this invention possesses excellent high-temperature grain size stability and resistance to anthill corrosion, as well as good processing performance and low cost, making it promising for applications in high-end fields such as heat exchange, new energy, and high-power computers.

[0012] The first aspect of this invention provides a corrosion-resistant and high-temperature resistant copper alloy tube, composed of P, X, copper, and unavoidable impurities, wherein X is selected from at least one of Fe, Mn, Cr, Ti, Zr, and Co, characterized in that, by mass percentage, 0.12%≤P≤0.36%, 0.0012%≤Fe≤0.50%, 0.0015%≤Mn≤0.45%, 0.0012%≤Cr≤0.40%, 0.0015%≤Ti≤0.30%, 0.0013%≤Zr≤0.27%, 0.0015%≤Co≤0.15%, and the total content of X satisfies 0.0012%≤X≤0.50%, the total amount of unavoidable impurities is ≤0.01%, and copper is the balance; The copper alloy tube comprises recrystallized grain structure G1, substructured grain structure G2, and deformed grain structure G3, with the following volume percentages: 0.50%≤G1≤40.0%, 55.0%≤G2≤95.0%, and 0.30%≤G3≤43.0%. G1 consists of grains with an orientation dispersion of ≤1°, G2 consists of grains with an orientation dispersion of 1° < ≤5°, and G3 consists of grains with an orientation dispersion of >5°.

[0013] This invention optimizes the alloy composition and synergistically controls the proportion of grain structure under different morphologies, thereby enabling the prepared copper alloy tube to have excellent high-temperature grain size stability and resistance to anthole corrosion.

[0014] This invention uses phosphorus (P) as one of the basic components. P reacts with Cu to form Cu-P phosphides, creating a continuous and dense Cu-P protective film on the substrate surface. This film effectively blocks the penetration of corrosive media and inhibits electrochemical corrosion. At the electronic structure level, the d-orbital electrons of P hybridize with the s- and p-orbital electrons of Cu, increasing the film resistivity. P has a strong affinity for oxygen (O), and during corrosion, it can form stronger covalent bonds with O in the environment, enhancing interatomic bonding and inhibiting stress-induced cracks or porosity during film growth, thus improving the film's density and stability. At the microstructure level, the Cu-P phase is distributed at grain boundaries and subgrain boundaries, effectively pinning grain boundaries and inhibiting abnormal grain growth under high-temperature service conditions. Simultaneously, the addition of P helps increase the recrystallization temperature of the alloy and slows down the recrystallization process.

[0015] In this invention, 0.12%≤P≤0.36% can be, for example, 0.12%, 0.15%, 0.175%, 0.20%, 0.225%, 0.25%, 0.275%, 0.30%, 0.325%, 0.35%, 0.36%, or any range consisting of any two endpoints.

[0016] When P < 0.12%, the Cu-P protective film on the substrate surface is insufficient, the film layer is discontinuous and porous, and a continuous and dense passivation layer cannot be formed. The resistance to acid and alkali and the resistance to ant hole corrosion are greatly reduced, and the alloy is prone to pitting corrosion and leakage failure. If the amount of Cu-P phosphide is too small, it cannot effectively pin the grain boundaries, and the grains grow abnormally at high temperature, which reduces the high temperature strength and creep resistance of the alloy. In addition, too low P content will lead to poor fluidity of liquid copper, and defects such as shrinkage cavities and porosity are prone to occur during casting and extrusion molding. When P > 0.36%, excessive P reacts with Cu to form a large amount of continuous network Cu3P brittle compounds, disrupting the continuity of the copper matrix and causing a sharp deterioration in the room temperature plasticity and hot and cold working properties of the alloy. Cracking is likely to occur during rolling and bending. Excessive Cu-P phosphides accumulate at grain boundaries, leading to a decrease in the thermal conductivity of the alloy, making it difficult to use in applications requiring high thermal conductivity. In addition, excessively high P content can easily lead to compositional segregation during smelting, with local phosphorus enrichment forming coarse brittle inclusions. During high-temperature service, the alloy is prone to intergranular cracking.

[0017] Preferably, 0.18% ≤ P ≤ 0.34%.

[0018] In this invention, the total content of X satisfies 0.0012%≤X≤0.50%, and can be, for example, 0.0012%, 0.0013%, 0.003%, 0.005%, 0.0075%, 0.01%, 0.0125%, 0.0325%, 0.05%, 0.0725%, 0.10%, 0.125%, 0.145%, 0.165%, 0.19%, 0.25%, 0.30%, 0.375%, 0.50%, or any range consisting of any two endpoints.

[0019] X can react with P to form high-melting-point metal phosphides, which enhance the pinning effect on grain boundaries and improve the high-temperature grain size stability of the alloy. At the same time, metal phosphides can participate in the formation of surface passivation film, forming a composite passivation film together with Cu-P protective film, and synergistically improving the corrosion resistance of the alloy.

[0020] When X < 0.0012%, the amount of metal phosphides in the system is insufficient, lacking effective grain boundary pinning phases and composite corrosion-resistant film components. The pinning effect is weakened, and the microstructure exhibits an excessively high proportion of G1, insufficient G2, a low proportion of special grain boundaries, and coarse grains in the finished product. At high temperatures, the grains are prone to coarsening, and the high-temperature resistance is significantly reduced. When X > 0.5%, the metal phosphides are excessive and coarsened, consuming a large amount of P and other alloying elements. Some grain boundaries lose their effective pinning effect, and the grain boundaries are occupied by brittle phosphides, forming a brittle network, leading to an increase in the proportion of G3 and a deterioration in the material's plasticity and processing performance. In addition, the excess alloying elements dissolve in the copper matrix, significantly weakening the alloy's thermal conductivity. At the same time, the mixture of multiple phosphides exacerbates the microgalvanic corrosion effect, accelerating the local corrosion rate.

[0021] In this invention, 0.0012%≤Fe≤0.50%, for example, can be 0.0012%, 0.002%, 0.003%, 0.005%, 0.0075%, 0.01%, 0.0125%, 0.0325%, 0.05%, 0.0725%, 0.10%, 0.125%, 0.135%, 0.15%, 0.20%, 0.35%, 0.48%, 0.50%, or any range consisting of any two endpoints.

[0022] In this invention, 0.0015%≤Mn≤0.45%, for example, can be 0.0015%, 0.0030%, 0.0050%, 0.0075%, 0.01%, 0.0125%, 0.0325%, 0.05%, 0.0725%, 0.10%, 0.125%, 0.135%, 0.15%, 0.25%, 0.375%, 0.43%, 0.45%, and any range consisting of any two endpoints.

[0023] When Fe or Mn reacts alone with P, it mainly forms Fe-P or Mn-P phosphides. These phosphides exhibit better thermal stability at brazing temperatures than Cu-P phosphides, effectively pinning grain boundaries and delaying grain growth. Some Fe or Mn dissolves in the Cu matrix, producing a solute dragging effect, further inhibiting grain boundary migration. In a anthill corrosion environment, Fe-containing compounds form on the alloy surface. 2+ or Mn 2+ The phosphate protective film is denser than the pure Cu-P film and can, to a certain extent, hinder the penetration of corrosive ions.

[0024] When Fe < 0.0012%, the amount of Fe-P phosphides is too small, the high-temperature grain boundary pinning effect fails, and abnormal grain growth is likely during high-temperature annealing or service, resulting in a decrease in the alloy's high-temperature strength and creep resistance. When Fe > 0.50%, the excess Fe, in addition to forming Fe-P phosphides, easily forms a Cu-Fe brittle phase with Cu. This coarse Cu-Fe brittle phase is distributed in a network along the grain boundaries, easily causing stress concentration during deformation and reducing the alloy's plasticity and processing properties. Furthermore, elemental Fe is a rust-prone element, inducing a micro-couple effect and reducing the alloy's corrosion resistance.

[0025] When Mn < 0.0015%, the Mn content is too low, resulting in weak deoxidation and desulfurization effects and increased matrix inclusions. The lack of Mn-P synergistic repair of passivation film defects also leads to insignificant corrosion resistance improvement. When Mn > 0.45%, excessive Mn dissolves in the Cu matrix, occupying the solid solution sites of other elements and reducing the alloy's plasticity, processing performance, and thermal conductivity. Simultaneously, Mn is prone to compositional segregation, leading to uneven composition in the smelted ingot.

[0026] In this invention, 0.0012%≤Cr≤0.40%, for example, can be 0.0012%, 0.0015%, 0.003%, 0.005%, 0.0075%, 0.01%, 0.0125%, 0.0325%, 0.05%, 0.0725%, 0.10%, 0.125%, 0.135%, 0.15%, 0.20%, 0.35%, 0.40%, or any range consisting of any two endpoints.

[0027] Cr partially dissolves in the matrix and partially forms Cr-P and Cu-Cr-P phosphides with P. Cr phosphides have high melting points and exhibit some pinning effect on grain boundaries. Cr promotes passivation of the alloy surface, forming a Cr-O / Cr-PO composite film. This film exhibits excellent stability in acidic media and effectively inhibits the initiation and propagation of pitting corrosion. When Cr < 0.0012%, the pinning effect of Cr-P phosphides on grain boundaries is insufficient, and the amount of Cu-Cr-P phosphides is also low, failing to form effective protection. When Cr > 0.40%, excessive Cr tends to aggregate at grain boundaries, forming a continuous brittle Cr-P network, reducing the alloy's ductility and significantly decreasing its thermal conductivity. Furthermore, high Cr content easily leads to compositional segregation, resulting in excessively high local hardness and decreased machinability.

[0028] In this invention, 0.0015%≤Ti≤0.30%, for example, can be 0.0015%, 0.003%, 0.005%, 0.0075%, 0.01%, 0.0125%, 0.0325%, 0.05%, 0.0725%, 0.10%, 0.125%, 0.135%, 0.15%, 0.28%, 0.30%, or any range consisting of any two endpoints.

[0029] In this invention, 0.0013%≤Zr≤0.27%, for example, can be 0.0013%, 0.0015%, 0.0030%, 0.0050%, 0.0075%, 0.01%, 0.0125%, 0.0325%, 0.05%, 0.0725%, 0.10%, 0.125%, 0.135%, 0.15%, 0.2%, 0.24%, 0.27%, and any range consisting of any two endpoints.

[0030] Both Ti and Zr readily form phosphides. Ti-P and Zr-P exhibit extremely high melting points and outstanding high-temperature grain size stability. At the high temperatures of brazing, Ti-P and Zr-P show almost no dissolution or coarsening, and their grain boundary pinning force is far stronger than Cu-P or Fe / Mn phosphides, effectively suppressing grain growth to the greatest extent. Regarding corrosion, Ti or Zr forms extremely stable Ti-O or Zr-O films on the surface, which, when combined with phosphate films, constitute a highly dense barrier layer.

[0031] When Ti < 0.0015%, high-melting-point Ti-P phosphides are scarce, significantly reducing the thermal stability of high-temperature grain size. Grains tend to coarsen under high-temperature conditions, and the deoxidation and impurity removal effects are insufficient, leading to an increase in matrix oxide inclusions. When Ti > 0.30%, excess Ti preferentially combines with O and P to form coarse Ti-P and Ti-O hard inclusions. These inclusions easily become crack initiation sites, reducing the mechanical and processing properties of the alloy. At the same time, hard particles cause stress concentration, inducing microcracks, and defects such as peeling and cracking are prone to occur during processing. Excessive Ti addition also increases raw material costs.

[0032] When Zr < 0.0013%, the Zr-P content is insufficient, resulting in weak grain boundary strengthening and failure of grain refinement. During high-temperature service, grain boundaries are prone to abnormal growth due to slippage. At the same time, the degassing and purification effect of Zr is insufficient, leading to an increase in porosity defects in the ingot. When Zr > 0.27%, excessive Zr generates large clusters of Zr-P inclusions. The hard agglomerated phase disrupts the continuity of the matrix, making the material prone to cracking during bending and stamping. Excessive Zr addition also increases raw material costs.

[0033] In this invention, 0.0015%≤Co≤0.15%, for example, can be 0.0015%, 0.003%, 0.005%, 0.0075%, 0.01%, 0.0125%, 0.0325%, 0.05%, 0.0725%, 0.10%, 0.125%, 0.135%, 0.15%, and any range consisting of any two endpoints.

[0034] Co exhibits moderate solid solubility in copper and a strong tendency for grain boundary segregation. Through the solute dragging effect, it can significantly increase the recrystallization temperature and delay grain boundary migration. With the synergy of P, Co can form Co-P and Cu-Co-P phosphides, forming an interwoven pinned network with Cu-P phosphides, exhibiting excellent stability at high brazing temperatures. Regarding resistance to termite corrosion, the Co released in the early stages of corrosion... 2+ Co(OH)2 or Co3(PO4)2 can be dynamically deposited in local alkaline areas to fill corrosion micropores and microcracks; the amorphous Co-P phosphide formed by Co and P has extremely high penetration blocking efficiency formate and chloride ions; at the same time, Co causes the alloy corrosion potential to shift positively to a moderate degree, thereby weakening the driving force of microcouple corrosion from the root.

[0035] When Co < 0.0015%, the insufficient Co content negates the synergistic benefit of the alloy's corrosion resistance and strength. When Co > 0.15%, excessive Co solution reduces the solid solubility of the copper substrate, causing the formation of other brittle phases, reducing the alloy's plasticity and impairing its thermal conductivity; furthermore, excessive Co addition is detrimental to cost control.

[0036] Preferably, 0.002% ≤ Co ≤ 0.13%.

[0037] In this invention, 0.50%≤G1≤40.0% can be, for example, 0.50%, 1.0%, 3.5%, 5.5%, 8.5%, 11.5%, 15.0%, 18.5%, 22.0%, 25.5%, 29.0%, 32.5%, 36.5%, 40.0%, or any range consisting of any two endpoints.

[0038] In this invention, 55.0%≤G2≤95.0% can be, for example, 55.0%, 58.5%, 61.5%, 65.0%, 68.5%, 72.0%, 78.5%, 81.0%, 85.5%, 89.0%, 92.5%, 95.0%, or any range consisting of any two endpoints.

[0039] In this invention, 0.30%≤G3≤43.0% can be, for example, 0.30%, 0.51%, 3.2%, 6.5%, 8.6%, 10.4%, 12.5%, 14.6%, 16.5%, 18.5%, 20.3%, 23.7%, 27.6%, 32.1%, 35.5%, 38.8%, 40.1%, 43.0%, or any range consisting of any two endpoints.

[0040] G1 plays a role in relieving stress and improving passivation; however, a volume fraction >40% can lead to uncontrolled high-temperature grain boundary migration. The substructure of G2 has a dual function: hindering G1 grain boundary migration and providing nucleation sites, forming the basis for grain size thermal stability and alloy corrosion resistance. When G2 <55.0%, the effect of hindering G1 grain boundary migration deteriorates, weakening high-temperature grain size stability. When G2 >95.0%, deformation storage energy increases, accelerating abnormal grain growth during high-temperature heat treatment. When G3 <43.0%, it provides effective phase nucleation sites; excessively high G3 concentrations increase recrystallization rates, leading to abnormal grain size growth at high temperatures. This invention synergistically controls the volume fractions of G1, G2, and G3 within the aforementioned ranges, balancing the recrystallization driving force and pinning resistance, thereby ensuring high-temperature grain stability after brazing and improving corrosion resistance.

[0041] In this invention, the average grain size of the copper alloy tube is 1.0~20.0μm, for example, it can be 1.0μm, 2.0μm, 4.0μm, 7.0μm, 8.5μm, 12.5μm, 13.5μm, 15.5μm, 16.9μm, 18.1μm, 19.2μm, 20.0μm, or any range consisting of any two endpoints.

[0042] While grain sizes <1.0 μm offer high strength, the excessively large grain boundary ratio means phosphides are insufficient to completely pin all grain boundaries, and the excessive interfacial energy reserve increases the tendency for grain coarsening at high temperatures. Simultaneously, excessively fine grains increase the corrosion area, potentially accelerating uniform corrosion. Grain sizes >20.0 μm result in a smaller initial grain boundary area, leading to further grain coarsening after brazing. In coarse-grained materials, pitting corrosion tends to form deeper and more concentrated pits. This invention controls the grain size within the aforementioned range, ensuring sufficient coverage of phosphide pinning points while reducing the continuous length of corrosion channels.

[0043] Preferably, the average grain size of the copper alloy tube is 2.5~6.8μm.

[0044] Preferably, the proportion of ∑3 grain boundaries satisfies 40.0%≤∑3≤70.0%, for example, it can be 40.0%, 42.5%, 46%, 48.5%, 50.5%, 53.1%, 56.5%, 58.7%, 60.5%, 63.5%, 65.6%, 68.2%, 70.0%, or any range consisting of any two endpoints.

[0045] The ∑3 special grain boundary exhibits low atomic arrangement distortion, with an interfacial energy only about one-tenth that of random large-angle grain boundaries. During high-temperature heat treatment, ∑3 can kinetically inhibit recrystallized grain growth. When the proportion of ∑3 grain boundaries is within the above range, it can significantly reduce the continuity of random large-angle grain boundaries, cut off the penetration channels of corrosive media along random grain boundaries, and improve corrosion resistance.

[0046] Preferably, the proportion of ∑9 grain boundaries satisfies 1.0%≤∑9≤8.0%, for example, it can be 1.0%, 1.5%, 2.3%, 2.5%, 3.0%, 3.5%, 4.0%, 4.5%, 5.0%, 5.5%, 6.0%, 6.5%, 7.2%, 8.0%, or any range consisting of any two endpoints.

[0047] ∑9 special grain boundaries possess low diffusion coefficients and interfacial activity. When the proportion of ∑9 grain boundaries is 1.0%~8.0%, their synergistic effect with ∑3 is enhanced: on the one hand, ∑9 grain boundaries can further disrupt the continuity of random grain boundaries, causing the corrosion path to be discontinuously distributed, hindering the continuous diffusion of corrosive ions, and improving corrosion resistance; on the other hand, the interaction between ∑9 and ∑3 grain boundaries can generate more low-energy grain boundary triple nodes, forming additional pinning on grain boundary migration and improving the high-temperature grain size stability.

[0048] Preferably, the mass ratio of P to X satisfies: 0.3 ≤ P / X ≤ 180.0.

[0049] The particle size distribution (P / X) determines the quantity, size, and distribution of metal phosphides. When P / X < 0.3, the X content in the alloy is relatively high, easily leading to the formation of excessive metal phosphides. These phosphides tend to aggregate and grow, forming coarse, brittle phases that can cause stress concentration during deformation, reducing the alloy's plasticity and processing performance. Simultaneously, excessive metal phosphides can create a galvanic cell effect, accelerating localized corrosion. When P / X > 180.0, the X content is relatively low, failing to generate sufficient metal phosphides to pin grain boundaries and form a composite passivation film, resulting in poor improvement in the alloy's high-temperature grain size stability and corrosion resistance.

[0050] Preferably, 0.4 ≤ P / X ≤ 177.0.

[0051] This invention preferably adds at least two elements from element X to enhance the corrosion resistance and high-temperature thermal stability of the alloy through the synergistic effect of multiple elements. At the microstructure control level, the phosphides generated by different alloying elements, due to their varying melting points, morphologies, and distributions, form a pinning system with a gradient structure, achieving multiple pinning of grain boundaries and subgrain boundaries and suppressing microstructure instability at high temperatures. At the surface protection level, the oxides and phosphides of each element synergistically construct a composite passivation film, which can complementaryly fill the structural defects of a single film layer, thereby enhancing the film's density and corrosion resistance.

[0052] Preferably, the X element comprises Fe and Mn, wherein the content is 0.0015%≤Fe≤0.06% and the Mn content is 0.0018%≤Mn≤0.13%.

[0053] Within the aforementioned elemental content range, Fe and Mn can react with P to form high-melting-point Fe-Mn-P composite phosphides that are not easily coarsened at brazing temperatures, exhibiting strong pinning force on grain boundaries. Simultaneously, Mn partially dissolves in the Cu matrix, creating a solute dragging effect that slows grain boundary migration. In corrosive environments, a Fe- and Mn-containing phosphate composite film can form on the alloy surface. The presence of Mn makes the oxide film denser, reducing pitting corrosion-sensitive areas and significantly inhibiting the expansion of anthill-like corrosion channels.

[0054] Preferably, the X element comprises Co, with 0.01% ≤ Co ≤ 0.13%, 0.0015% ≤ [Fe, Mn, Cr, Ti, Zr] ≤ 0.06%, where [Fe, Mn, Cr, Ti, Zr] represents at least one selected from Fe, Mn, Cr, Ti, and Zr. 0.0015% ≤ Fe ≤ 0.05%, 0.0015% ≤ Mn ≤ 0.025%, 0.0015% ≤ Cr ≤ 0.026%, 0.0015% ≤ Ti ≤ 0.032%, and 0.0015% ≤ Zr ≤ 0.027%.

[0055] When Co coexists with [Fe, Mn, Cr, Ti, Zr], due to the similarity of Co's atomic radius to Cu, some Co dissolves in the copper matrix, increasing the recrystallization temperature. The remaining Co forms [Fe, Mn, Cr, Ti, Zr]-Co-P type composite phosphides. These composite phosphides constitute a thermally stable framework, stabilizing the fine-grained structure and improving high-temperature grain size stability. Simultaneously, Co fills vacancies in the [Fe, Mn, Cr, Ti, Zr] passivated matrix, forming an amorphous, multi-component, dense composite phosphide layer that hinders the penetration of corrosive ions.

[0056] More preferably, the X element comprises Co and Zr, with 0.05%≤Co≤0.13% and 0.0015%≤Zr≤0.027%.

[0057] Preferably, X comprises Co, Fe and Mn, wherein 0.01%≤Co≤0.13%, 0.0015%≤Fe≤0.012%, and 0.0015%≤Mn≤0.022%.

[0058] Within the aforementioned elemental content range, Fe-Mn-Co-P multi-component phosphides will form. The high mixing entropy of the multi-element formation lowers the Gibbs free energy of the phosphides, significantly improving their thermodynamic stability. They remain almost undissolved and do not coarsen during brazing at 800–1000℃, exhibiting persistent grain boundary pinning. The solid-solution Mn and Co generate strong combined solute drag, significantly inhibiting static recrystallization grain growth. The corrosion film is a (Fe-Mn-Co)-PO4 composite salt, possessing density, self-healing ability, and a wide pH stability range, further reducing the depth of anthill corrosion.

[0059] Preferably, X comprises Co, Ti and Zr, wherein 0.01%≤Co≤0.13%, 0.0015%≤Ti≤0.012%, and 0.0015%≤Zr≤0.011%.

[0060] Ti and Zr are strong phosphide-forming elements, preferentially forming Ti-Zr-P phosphides, which have a much stronger pinning force on grain boundaries than intermetallic compounds. Co partially dissolves into the Ti-Zr-P lattice or aggregates at the interface, reducing interfacial energy and preventing particle coarsening. The synergistic effect of Ti, Zr, and Co further reduces the grain size after brazing, and the resulting composite protective film further improves the alloy's resistance to anthill corrosion.

[0061] Preferably, the X element comprises Co, Cr, Ti and Zr, with 0.08%≤Co≤0.13%, 0.0015%≤Cr≤0.006%, 0.0015%≤Ti≤0.011%, and 0.0015%≤Zr≤0.007%.

[0062] Preferably, the X element comprises Co, Mn, Cr, Ti and Zr, wherein 0.09%≤Co≤0.13%, 0.0015%≤Mn≤0.011%, 0.0015%≤Cr≤0.005%, 0.0015%≤Ti≤0.009%, and 0.0015%≤Zr≤0.0065%.

[0063] Preferably, the X element comprises Co, Fe, Mn, Cr, Ti, and Zr, wherein 0.10%≤Co≤0.13%, 0.0015%≤Fe≤0.008%, 0.0015%≤Mn≤0.0065%, 0.0015%≤Cr≤0.005%, 0.0015%≤Ti≤0.0068%, and 0.0015%≤Zr≤0.004%.

[0064] When X is composed of 4 or more elements such as Cr, Ti, Zr, and Co (including Fe and / or Mn), the mixing entropy is greater, the thermal stability of the grain size is further improved, and the grain size is more stable after brazing. The corrosion product film is an amorphous dense film containing polyphosphates and oxides, which hinders the grain boundary diffusion channels and significantly improves the resistance to anthole corrosion.

[0065] In this invention, unavoidable impurities include elements such as S, As, C, H, and O, with a total mass percentage of <0.01%. Among these unavoidable impurities, S and As readily form low-melting-point brittle compounds with copper, distributed along grain boundaries, leading to hot brittleness and reducing the material's machinability; their content percentages are both ≤0.0005%. When C exists in copper as inclusions such as carbides, it reduces the alloy's plasticity, toughness, and machinability, becoming a stress concentration source and easily initiating cracks. Furthermore, excessive C deteriorates the material's surface quality, increasing the risk of corrosion cracking; its content percentage is ≤0.0015%. During copper smelting, if the oxygen content of the melt is too high, it reacts with hydrogen to generate water vapor, forming bubbles during solidification, leading to defects such as porosity and looseness in the castings, and in severe cases, rendering the copper alloy unusable. In addition, excessive oxygen content also leads to oxide inclusions, reducing the alloy's mechanical properties, corrosion resistance, and machinability; therefore, the content percentages of H and O are both ≤0.001%.

[0066] The copper content is the balance, which is 100% minus the contents of P, X and all impurities.

[0067] Preferably, the ratio of the sum of the volume percentages of G1 and G2 to the volume percentage of G3 satisfies: (G1+G2) / G3≥1.5.

[0068] (G1+G2) / G3 represents the ratio of "stable" structure that has completed recovery or recrystallization to residual highly distorted structure in the material. When the ratio is too low, G3 accounts for a high proportion, and a large amount of deformation stored energy is released instantaneously during brazing heating, triggering explosive nucleation and rapid growth of recrystallized grains, resulting in severe grain coarsening and a sharp increase in grain size at high temperatures; at the same time, the dislocation entanglement regions of the deformed structure are preferential active sites for corrosion, and the depth of anthill corrosion increases significantly.

[0069] Preferably, the volume ratio of G2 to G1 satisfies: 1.5≤G2 / G1≤65.0.

[0070] The G2 / G1 ratio reflects the proportion of substructures relative to recrystallized grains. A low G2 / G1 ratio indicates insufficient substructure, weak subgrain boundary pinning, and easy grain growth at high temperatures. Simultaneously, a low G1 ratio results in fewer internal defects within the grains, hindering the formation of a continuous and dense corrosion product film and reducing corrosion resistance. Conversely, a high G2 / G1 ratio, while providing strong pinning, leads to insufficient material plasticity due to the low recrystallization ratio, making processing difficult. Furthermore, excessive subgrain boundaries can become rapid diffusion channels for corrosion, increasing susceptibility to anthill corrosion. Maintaining a G2 / G1 ratio between 1.5 and 65.0 balances pinning strengthening and plasticity, reducing grain growth rates at high brazing temperatures, resulting in a suitable corrosion film substrate and superior corrosion resistance.

[0071] Preferably, in the copper alloy tube, 46.0%≤∑3+∑9≤70.0%.

[0072] Low-energy special grain boundaries ∑3 and ∑9 can effectively segment grains, increase interface density, hinder dislocation movement and grain boundary migration, improve strength, and refine grains. Simultaneously, the coherent ∑3 interface has extremely low energy, is insensitive to corrosive media, and can interrupt continuous, random, large-angle grain boundary corrosion pathways. When ∑3 + ∑9 < 46.0%, insufficient special grain boundary segmentation results in a small contribution to grain refinement, and corrosion easily propagates continuously along large-angle grain boundaries; when ∑3 + ∑9 > 70.0%, excessive special grain boundary intersections lead to localized stress concentration, easily inducing a micro-cell effect in special grain boundary corrosion.

[0073] Preferably, in the copper alloy tube, 8.0 ≤ ∑3 / ∑9 ≤ 35.0.

[0074] When ∑3 / ∑9 < 8.0, the energy and activity of the incoherent interface are high, resulting in decreased corrosion resistance; when ∑3 / ∑9 > 35.0, the proportion of ∑3 grain boundaries is too high, sacrificing the additional strengthening and pinning effect brought by ∑9, and the overall performance is not optimal.

[0075] Preferably, 8.0 ≤ ∑3 / ∑9 ≤ 15.0.

[0076] By maintaining the values ​​of ∑3+∑9 and ∑3 / ∑9 within the above range, this invention can maximize the dual benefits of special grain boundaries for grain refinement and corrosion resistance.

[0077] In this invention, the average grain size of the copper alloy tube after heat treatment at 1000℃ for 30 min is 195~420 μm, and the maximum corrosion depth after alternating hot and cold corrosion in a 0.5% volume fraction formic acid aqueous solution atmosphere for 8 days is 40~78 μm.

[0078] The second aspect of the present invention provides a method for preparing the above-mentioned corrosion-resistant and high-temperature resistant copper alloy tube, comprising the following steps: batching and smelting → horizontal continuous casting → planetary rolling → stretching → finished product annealing.

[0079] In the batching and smelting process, the smelting temperature is controlled at 1100~1230℃. Charcoal is used to cover the surface of the molten copper. Electromagnetic stirring is continuously activated throughout the smelting process, with a current range of 85~110A and a frequency range of 25~50Hz. Traditional feeding methods typically involve simultaneous feeding or sequential feeding according to melting point. While this invention can also utilize traditional feeding methods, it preferentially employs a reverse feeding method based on reactivity priority. S1: Add copper raw materials. When the melt temperature reaches 1150~1230℃, add Cu-P master alloy according to the ratio and keep it at the temperature for 10~15 minutes to make P evenly distributed in the copper liquid. S2: Reduce the melt temperature to 1100~1140℃, and selectively add intermediate alloys containing Co, Fe, Mn, Cr, Ti and Zr in order of increasing affinity with P. After each intermediate alloy is added, hold the temperature for 5~10 minutes to ensure uniform composition distribution.

[0080] Traditional processes typically add intermediate alloys simultaneously or sequentially according to their melting points. This often leads to a rapid reaction between strong phosphides and phosphorus (P) in the melt, forming coarse primary phosphides that are difficult to break down during subsequent processing, have low density, and poor pinning effect. In contrast, this invention first adds a Cu-P intermediate alloy, allowing P to be preferentially and uniformly distributed in the copper melt, creating a melt environment with high P chemical potential. Weak affinity elements (Co, Fe, Mn, Cr) are then added; some are dissolved in the copper matrix, while others exist as solute atoms. These elements can form compounds with P or, after the addition of Ti and Zr, diffuse onto the surface of newly formed Ti-Zr-P nanocores, forming composite phosphides in situ. Ti and Zr are added last to ensure they react with P at the lowest possible temperature, reducing the nucleation driving force.

[0081] In the horizontal continuous casting process, the casting temperature is 1145~1175℃, the traction speed is 275~425mm / min, the cooling water inlet flow rate is 15~55L / min, and the inlet and outlet water temperature difference is 15~25℃.

[0082] Preferably, in this invention, the traction speed is inversely proportional to the casting temperature. When the casting temperature is close to the upper limit (1170~1175℃), the traction speed is at the lower limit (275~320mm / min); when the casting temperature is close to the lower limit (1145~1150℃), the traction speed is at the upper limit (380~425mm / min).

[0083] The traction speed is inversely linked to the casting temperature, achieving a dynamic balance between the heat dissipation rate at the solidification front and the compensation for solidification shrinkage. When the casting temperature is too high, the traction speed is reduced to allow more time for heat dissipation at the solidification front, preventing overheating in the central region and resulting in grain coarsening. When the temperature is too low, the traction speed is increased to prevent excessive cooling, which could lead to an excessively thick surface quench layer and uneven microstructure in the core, ultimately resulting in an ingot with a uniform cross-sectional microstructure and fine grains. This microstructure is inherited during planetary rolling, resulting in more uniform deformation flow lines, providing a homogeneous foundation for obtaining a suitable G1, G2, and G3 ratio during subsequent annealing.

[0084] In the planetary rolling process, the feed rate is 1.5~2.4 m / s, the discharge rate is 15~25 m / s, and the deformation is 90~92%. Preferably, the ingot is preheated to 350~420℃ before rolling to give the microstructure a certain degree of adaptability to thermal deformation. After rolling, the billet is rapidly cooled to below 100℃ within 0.5~1.5 seconds, followed by natural air cooling.

[0085] This invention employs a large deformation margin of 90-92%, utilizing the deformation temperature rise effect to push the rolling temperature to 650-850℃, fully triggering dynamic recrystallization. This refines the coarse as-cast microstructure into uniform recrystallized grains, eliminating the inheritance of the casting microstructure. After rolling, cooling is performed, rapidly reducing the billet temperature from 650-850℃ to below 100℃ within 0.5-1.5 seconds. This quickly "freezes" the microstructure just as recrystallization is complete and before the grains have grown, controlling further growth of the recrystallized grains while preserving some metastable interface structures formed due to incomplete grain boundary migration. Furthermore, rapid cooling keeps phosphorus and other alloying elements in a supersaturated solid solution state, preventing premature coarsening of phosphides at high temperatures and providing sufficient supersaturated solute reserves for uniform phosphide formation during subsequent annealing.

[0086] Preferably, the stretching can be done by straight stretching, continuous stretching or disc stretching, and the stretching passes are at least 4 passes. The elongation coefficient of the first 2 passes is 1.35~1.65; the elongation coefficient of the subsequent passes is 1.15~1.30, and the pipe temperature is controlled not to exceed 50°C.

[0087] This invention achieves sufficient and uniform plastic deformation across the entire tube cross-section through the first two large-scale elongation deformation passes, significantly increasing dislocation density and smoothing out the deformation difference between the surface and core through strong rheological processes. Subsequent small-scale elongation deformation passes avoid introducing excessive deformation, preventing a significant increase in dislocation density. By controlling the elongation coefficients of different passes, this invention retains sufficient deformation storage energy to provide a suitable recrystallization driving force for subsequent annealing, while also forming an appropriate number of substructure dislocation cells. This provides a structural basis for the rapid formation of a uniform G2 substructure in the recovery section of the finished product's annealing process. Maintaining the tube temperature below 50°C effectively avoids excessive recovery caused by artificial aging or high temperatures in traditional processes, thus maintaining deformation storage energy within the optimal window.

[0088] In the finished product annealing process, gradient annealing is adopted, which includes recovery stage annealing and recrystallization stage annealing in sequence. The annealing conditions for the recovery stage are as follows: in an inert atmosphere, the pipe is kept at 260~290℃ for 40~80min; The annealing conditions for the recrystallization stage are as follows: in an inert atmosphere, the pipe is kept at 570~595℃ for 30~50min; After the gradient annealing is completed, the furnace is cooled to room temperature.

[0089] The finished product annealing of this invention adopts a two-stage gradient temperature control system of recovery and recrystallization, which is different from traditional isothermal annealing. Its core mechanism is: low-temperature recovery to build a uniform substructure matrix, and high-temperature recrystallization to regulate the distribution of grain boundary characteristics, thereby achieving synergistic optimization of G1, G2, G3 grain structure and ∑3, ∑9 special grain boundary distribution.

[0090] During the recovery stage, atomic diffusion is restricted, dislocation rearrangement occurs in the deformed structure, forming low-energy subgrain boundaries. This divides the grain interior into a uniform subgrain structure G2, while simultaneously inhibiting large-angle grain boundary migration and recrystallization nucleation. This stage transforms some of the deformed structure G3 into G2, providing uniform nucleation sites for subsequent recrystallization. At temperatures below 260℃ or times shorter than 40 minutes, recovery is insufficient, resulting in more G3 residue, increased residual stress, and a heightened risk of mixed crystal formation, which impairs corrosion resistance. At temperatures above 290℃ or times longer than 80 minutes, excessive dislocation annihilation occurs, leading to insufficient recrystallization driving force and deterioration of high-temperature stability.

[0091] It should be noted that the optimal process parameters for the recovery stage are adjusted according to the alloy composition: when X is a quaternary or higher system, the upper limit temperature (280~290℃) is used and the holding time is shortened (40~60min); when X is a ternary or lower system, the lower limit temperature (260~280℃) is used and the holding time is extended (60~80min) to ensure sufficient recovery without excessive consumption of stored energy. At the end of the recovery stage, the microstructure is a mixed state of "G3+G2+G1", with a significant increase in the proportion of G2.

[0092] During the recrystallization stage, sufficient atomic diffusion occurs, and uniform subgrains migrate through subgrain boundaries to become recrystallization nuclei, promoting the continuous transformation of G3 to G1. The microstructure is dominated by ∑3 special grain boundaries, with a small amount of ∑9 associated with it. The key to this process is to control the degree of recovery to fully consume the deformed microstructure while avoiding excessive reduction of G2 due to complete recrystallization; and to inhibit grain growth through the synergistic effect of phosphide pinning and solute dragging. ∑3 grain boundaries divide grains, pin dislocations, and break the random grain boundary network, blocking corrosion channels and improving corrosion resistance. An appropriate amount of ∑9 further enhances the corrosion resistance of ∑3. If the temperature is below 570℃ or the time is less than 30 min, recrystallization is insufficient, the amount of special grain boundaries generated is small, and the optimization of grain boundary characteristic distribution is inadequate; if the temperature is too high or the time is too long, the grains coarsen, the ∑3 / ∑9 ratio deviates from the suitable range, the continuity of random grain boundaries is restored, and the corrosion resistance and high-temperature stability deteriorate.

[0093] The recrystallization process employs furnace cooling, which offers three key benefits: first, it promotes the segregation of microalloying elements towards the interface, reducing interfacial energy; second, it drives phosphides to disperse along grain boundaries, enhancing high-temperature pinning; and third, it eliminates thermal stress, preventing the introduction of dislocations and distortions, and ensuring a low-defect microstructure. The synergistic effect of these three benefits provides reliable microstructural protection for the alloy brazing process in high-temperature service and in environments prone to anthill corrosion.

[0094] If it is necessary to prepare internally threaded tubes, online annealing and internal thread forming can be performed sequentially after stretching and before the finished product is annealed.

[0095] In the online annealing process, the online annealing temperature is 530~850℃, the feed rate is 330~430m / min, and the online annealing temperature T (℃) and feed rate v (m / min) are controlled to satisfy the following relationship: T × v = 2.0×10 5 ~ 2.8×10 5 ℃·m / min.

[0096] The product of temperature and rate (temperature-rate product) represents the effective thermal energy absorbed per unit mass of material in the induction heating zone. Controlling this product within the specified window ensures similar properties for products of different specifications after annealing, significantly improving batch-to-batch microstructure consistency. After online annealing, the material undergoes complete recrystallization, eliminating work hardening and forming fine, uniform equiaxed grains, while simultaneously generating a large number of Σ3 low-energy special grain boundaries. Σ3 grain boundaries exhibit excellent genetic stability in subsequent internal thread forming and finished product annealing; they do not completely disappear due to cold deformation but are retained as "grain boundary characteristic seeds," reactivated and proliferating in the finished product annealing stage, thereby regulating the ratio of Σ3+Σ9. When the temperature-rate product is too low, recrystallization is insufficient, leaving residual deformed structures, leading to increased strength and decreased corrosion resistance; when it is too high, grains coarsen rapidly, the proportion of special grain boundaries decreases, and grains easily grow at high temperatures.

[0097] Preferably, in the internal thread forming process, the forming rate is 25~55m / min, and the equivalent cold working deformation (calculated by the change in cross-sectional area) is 12~18%.

[0098] By controlling the equivalent cold working deformation and forming rate within the above range, a large number of stacking faults and special grain boundary precursors can be accumulated in the grains, providing nucleation sites for Σ3 special grain boundaries in the subsequent annealing of the finished product, thereby stabilizing the high-temperature grain size and improving the corrosion resistance of the alloy.

[0099] Before planetary rolling, the ingots obtained from horizontal continuous casting can be milled to uniformly remove the surface oxide layer and casting defects, ensuring the surface quality of subsequent rolling. This is a conventional technique and will not be described in detail here.

[0100] The third aspect of this invention provides applications of the aforementioned copper alloy tubes or copper alloy tubes prepared by the aforementioned method in heat exchange, new energy, high-power computers, and other fields. The copper alloys of this invention possess excellent high-temperature grain size stability and corrosion resistance, and have good processing performance. They can be widely used in heat exchange, new energy, high-power computers, and other fields, and are particularly suitable for products such as heat exchange tubes, condensers, and evaporators that require brazing and operate in corrosive environments. Attached Figure Description

[0101] Figure 1 The grain orientation dispersion diagrams are for Examples 1, 21, 50, 58, 61, 64, 69, 70 and Comparative Examples 1 and 3.

[0102] Figure 2 The distribution diagrams show the low-energy special grain boundaries ∑3 and ∑9 of Examples 1, 21, 50, 58, 61, 64, 69, 70 and Comparative Examples 1 and 3.

[0103] Figure 3 Metallographic images of grain size after heat treatment at 1000℃ for 30 min for Examples 1, 21, 50, 58, 61, 64, 69, 70 and Comparative Examples 1 and 3.

[0104] Figure 4 Metallographic images of the anthill corrosion depth after 8 days of corrosion in Examples 1, 21, 50, 58, 61, 64, 69, 70 and Comparative Examples 1 and 3 in a 0.5% formic acid concentration environment.

[0105] Figure 5 The image shows the macroscopic morphology of cracking in the alloys of Comparative Examples 4-9 during planetary rolling. Detailed Implementation

[0106] To make the objectives, technical solutions, and advantages of the embodiments of this application clearer, the technical solutions of the embodiments of this application will be clearly and completely described below with reference to the data and accompanying drawings. Obviously, the described embodiments are only a part of the embodiments of this application, and not all of them. All other embodiments obtained by those skilled in the art based on the embodiments of this application without creative effort are within the scope of protection of this application.

[0107] The microstructure and performance parameters involved in the embodiments and comparative examples of this invention were determined using the following methods.

[0108] High-Temperature Grain Size Test: The finished copper tube was placed in a tube furnace at 1000℃ and held for 30 minutes before being removed and cooled to room temperature with water. The high-temperature treated copper tube was then prepared into metallographic samples, mechanically polished until the surface was free of obvious scratches, and subsequently etched. The etching solution was prepared by mixing ammonium persulfate (2g) and deionized water (20mL), placed in a plastic beaker, and the sample was immersed in the etching solution with gentle agitation. After etching for 20 seconds, the sample was rinsed with water, dehydrated with ethanol, and finally dried with a hair dryer. The metallographic structure was observed using a ZEISS Axio Lab.A1 optical microscope, and the average grain size was measured according to GB / T 6394-2017 "Method for Determination of Average Grain Size of Metals". For each example and comparative example, five parallel samples were taken circumferentially along the cross-section for measurement, and the average value was taken as the final result.

[0109] Anthole Corrosion Resistance Test: The anthole corrosion resistance test was conducted using the apparatus disclosed in patent CN110546286B. The test method is as follows: Before the test, the alloy sample was sequentially subjected to ultrasonic cleaning with ethanol for 3 minutes, ultrasonic degreasing with acetone for 3 minutes, and ultrasonic cleaning with deionized water for 3 minutes, and finally dried to ensure the surface was clean and free of oil. The pretreated sample was placed in a simulated anthole corrosion environment with a 0.5% (v / v) formic acid solution as the corrosive medium. A water bath circulation temperature control was used, alternating between constant temperature exposure at 55℃ for 24 hours and constant temperature exposure at 25℃ for 24 hours, for a total corrosion cycle of 8 days. The sample length was 10 cm, with 9 cm of the exposed portion inside the container. The sample did not come into contact with the formic acid solution, and the bottom of the sample inside the container was approximately 5 cm from the top surface of the formic acid solution. After the experiment, metallographic samples were transversely cut at the midpoint of the length of each copper tube. The cross-sections were mechanically polished until the surface was free of obvious scratches. Using a metallographic microscope, the maximum depth of the anthill corrosion was observed and recorded from the outer surface of the copper tube inward along the radial direction of the cross-section. The average of the maximum corrosion depths of 15 parallel samples was taken as the characterizing maximum corrosion depth of the alloy.

[0110] Grain boundary, average grain size, and grain structure testing methods: EBSD was used to study the grain size and the proportion of low-energy special grain boundary angles in the alloy. The experimental equipment used was a ZEISS Gemini 2 equipped with an Oxford Symmetry EBSD detector. The test parameters were set as follows: accelerating voltage 20kV, working distance 15-20mm, step size 0.25μm, and detection area 49703.4μm per image. 2 (260.5μm × 190.8μm). Before testing, the sample underwent mechanical polishing and argon ion beam polishing to remove the residual stress layer on the surface. Data processing was performed using AZTEC software. The definition and statistical basis of grain boundary types are as follows: grain boundaries with an orientation difference angle ≥ 15° are defined as large-angle grain boundaries, and grain boundaries with an orientation difference angle ≤ 2° < 15° are defined as small-angle grain boundaries; low-energy special grain boundaries are identified according to the Brandon criterion, and the orientation difference relationship of Σ3 grain boundaries is... <111> Characterized by a 60° axial rotation, the orientation difference relationship of the Σ9 grain boundaries is... <110> It is characterized by a 38.94° rotation of the axis.

[0111] The proportions of Σ3 and Σ9 grain boundaries are calculated using the "length ratio method," and the specific calculation formula is as follows: Σ3 grain boundary ratio = L Σ3 / L 总 ×100% Σ9 grain boundary ratio = L Σ9 / L 总 ×100% The total proportion of Σ3 and Σ9 grain boundaries = (L Σ3 +L Σ9 ) / L total × 100% Among them, L Σ3 To determine the total length of all Σ3 grain boundaries within the detection region, L Σ9 To determine the total length of all Σ9 grain boundaries within the detection region, L 总 The total length of all large-angle grain boundaries (≥15°) within the detection area is measured in μm. All three values ​​were obtained statistically using AZTEC software.

[0112] The average grain size of the product was statistically analyzed using the equivalent circle diameter method. This method converts the grain area into the diameter of a circle with an equal area as the grain size measure. The arithmetic mean of the equivalent circle diameters of all grains within the statistical region was then taken to obtain the average grain size. Three different fields of view were randomly selected for each sample for statistical analysis, and the average value was taken as the final data.

[0113] GOS (Grain Orientation Spread) refers to the average angle of deviation of all points within a grain from the average orientation of that grain. In EBSD data, it can be statistically derived using AZTEC software. The criteria for recrystallized grain structure G1, substructured grain structure G2, and deformed grain structure G3 based on GOS classification are as follows: G1 (Recrystallized): A newly nucleated, distortion-free grain structure with an internal dislocation density close to that of the annealed state, completely surrounded by large-angle grain boundaries, and an internal GOS ≤ 1°.

[0114] G2 (Substructured): The original deformed grains are restored, and dislocations rearrange to form clear subgrains, but no large-angle grain boundary migration or new grain nucleation occurs. The internal grain angle is 1° < GOS ≤ 5°.

[0115] G3 (Deformed): No significant recovery occurred after cold deformation. The grains were elongated / distorted, dislocations were tangled, there was no clear subgrain structure, and the internal GOS of the grains was >5°.

[0116] The grain size distribution of G1, G2, and G3 is calculated using the "area ratio method." In a three-dimensional structure, the area ratio can be approximately equivalent to the volume ratio. The specific formulas for calculating the volume ratio of the grain size distribution of G1, G2, and G3 are as follows: G1 volume percentage = A G1 / A 总 ×100% G2 volume percentage = A G2 / A 总 ×100% G3 volume percentage = A G3 / A 总 ×100% Among them, A G1 To detect the total area of ​​all G1s within the detection region, A G2 To detect the total area of ​​all G2s within the detection region, A G3 To detect the total area of ​​all G3s within the detection region, A 总 The area is the total area of ​​the EBSD effective calibration area within the detection area, with the area unit being μm². All three values ​​were obtained statistically using AZTEC software.

[0117] Processing performance evaluation: During the planetary rolling stage of the sample, the surface quality of the sample after planetary rolling is observed by the naked eye. If the surface is smooth and there are no visible cracks, the sample has good processing performance and can be judged as qualified. If the surface is rough and there are obvious cracks, the sample is marked as a cracked sample and the processing performance of the sample is unqualified.

[0118] Comparative Examples 1-3 and Example 70 were prepared according to the following steps: S1, Ingredient Batching and Melting Prepare electrolytic copper and Cu-P (P content 14%), Cu-Co (Co content 10%), Cu-Fe (Fe content 10%), Cu-Mn (Mn content 10%), Cu-Cr (Cr content 10%), Cu-Ti (Ti content 10%), and Cu-Zr (Zr content 5%) master alloys according to the target composition. All master alloy contents refer to mass percentages.

[0119] First, control the melting temperature at 1160℃, cover the surface of the copper liquid with charcoal to prevent oxidation, add Cu-P master alloy, and hold for 15 minutes; then control the temperature at 1120℃, and selectively add Cu-Co, Cu-Fe, Cu-Mn, Cu-Cr, Cu-Ti, and Cu-Zr master alloys in sequence according to the target composition, holding for 8 minutes after each master alloy is added.

[0120] During the smelting process, the electromagnetic stirring was kept on continuously, with a current of 96A and a frequency of 35Hz.

[0121] S2, Horizontal Continuous Casting

[0122] The casting temperature was controlled at 1163℃, the traction speed was set at 351mm / min, the cooling water inlet flow rate was 35L / min, and the inlet and outlet water temperature difference was controlled at 18℃. After continuous casting, ingots with consistent specifications were obtained.

[0123] S3, Milling Surface

[0124] After straightening, the ingot is milled to uniformly remove the surface oxide layer and casting defects, ensuring the surface quality of subsequent rolling.

[0125] S4, Planetary Rolling

[0126] The milled ingot is preheated to 380℃, the feed rate is controlled at 2.1m / s, the discharge rate is controlled at 21m / s, and the rolling deformation is stabilized at 90%. After rolling, laminar cooling is performed, and the billet is rapidly cooled from 700~830℃ to below 100℃ within 0.5~1.5 seconds, followed by natural air cooling.

[0127] S5, Panla

[0128] A total of four coiling processes are set up. The elongation coefficient of the first two processes is controlled at 1.6, and the elongation coefficient of the last two processes is controlled at 1.3. During the coiling process, the pipe temperature is controlled not to exceed 50℃.

[0129] S6, Finished product annealing

[0130] The sample after being drawn into a disc was annealed in a nitrogen atmosphere at a temperature of 540℃ for 50 minutes. After annealing, it was air-cooled to room temperature to obtain a finished copper alloy tube with an outer diameter of 7 mm and a wall thickness of 0.7 mm.

[0131] The difference between the preparation process of Examples 1-69 and Example 70 is that the finished product is annealed using gradient annealing. The annealing temperature in the recovery stage is 280℃ and the holding time is 60min; the annealing temperature in the recrystallization stage is 575℃ and the holding time is 35min. After the gradient annealing is completed, the product is cooled to room temperature in the furnace.

[0132] The P content of Comparative Example 4 was 0.40%, which exceeded the upper limit of the present invention. It cracked during the planetary rolling process and could not be processed further.

[0133] The total content of alloying element X in Comparative Examples 5-9 exceeded the upper limit of this invention, and cracking occurred in the planetary rolling process, making subsequent processing impossible.

[0134] The specific components of the embodiments and comparative examples of this invention are shown in Table 1, and the test results are shown in Tables 2-3, respectively. Figures 1-4 .

[0135] In Examples 1-4, with increasing Fe content, the volume percentage of G1 gradually increased (from 1.0% to 15.5%), the percentage of G3 gradually decreased (from 40.0% to 1.1%), while the percentage of G2 showed a trend of first increasing significantly and then gradually decreasing (from 59.0% in Example 1 to 89.7% in Example 2, and then decreasing to 83.4%), (G1+G2) / G3 increased from 1.5 to 89.9, and G2 / G1 decreased from 59.0 to 5.4. The initial increase and subsequent decrease in the percentage of G2 is attributed to the competitive regulation of recovery and recrystallization by solute content. At low solute contents, solute atoms drag dislocation movement and grain boundary migration inhibit recrystallization nucleation, and deformed grains preferentially form substructures through recovery, thus increasing the percentage of G2. With further increases in solute content, dislocation energy accumulation increases, and the recrystallization driving force strengthens to exceed the critical pinning resistance corresponding to substructure stability. Substructure grains transform into recrystallized grains through grain boundary migration, and the percentage of G2 decreases accordingly. With increasing Fe content, the proportions of ∑3 and ∑9 increase, with ∑3+∑9 rising from 50.9% to 67.3%, and ∑3 / ∑9 decreasing from 32.9 to 11.0. The finished grain size refines from 6.3 μm to 4.0 μm. After heat treatment at 1000℃ for 30 min, the grain size decreases from 412.3 μm to 286.1 μm (a decrease of 30.6%), and the maximum corrosion depth after 8 days of 0.5% formic acid etching decreases from 77.2 μm to 74.5 μm (a decrease of 3.5%). This indicates that increasing Fe content improves the substructure and optimizes the distribution of grain boundary characteristics, thus significantly inhibiting high-temperature grain growth. However, the Fe-based phosphide / oxide film lacks sufficient density, limiting its ability to block the expansion of corrosion channels.

[0136] In Examples 5-29, as the contents of Mn, Cr, Ti, Zr, and Co increased, the volume fractions of G1, G2, G3, ∑3, and ∑9 showed the same trend as Fe, but the proportions of different structural grain structures and special grain boundaries changed significantly with different elements.

[0137] In Examples 5-8, with increasing Mn content, (G1+G2) / G3 increased from 2.0 to 110.1, G2 / G1 decreased from 54.6 to 4.4, ∑3+∑9 increased from 52.5% to 67.4%, and ∑3 / ∑9 decreased from 29.9 to 10.4, resulting in a refined grain size from 6.2 μm to 3.8 μm. After heat treatment at 1000℃ for 30 min, the grain size decreased from 405.6 μm to 292.0 μm (a decrease of 28.0%), and the maximum corrosion depth after 8 days of 0.5% formic acid corrosion decreased from 76.5 μm to 60.5 μm (a decrease of 20.9%). The increased Mn content also promoted the transformation of deformed structures into substructures and recrystallized states, and increased the proportion of special grain boundaries, thereby strengthening the high-temperature resistance to grain coarsening. Simultaneously, the dense phosphate film formed by Mn in the corrosive medium, combined with the optimized grain boundary characteristic distribution, synergistically reduced the corrosion depth.

[0138] In Examples 9-12, increasing the Cr content increased the (G1+G2) / G3 ratio from 2.5 to 99.0, decreased the G2 / G1 ratio from 50.0 to 2.7, increased the ∑3+∑9 ratio from 53.0% to 66.9%, and decreased the ∑3 / ∑9 ratio from 28.4 to 11.6, resulting in a refinement of the finished grain size from 6.0 μm to 3.6 μm. After heat treatment at 1000℃ for 30 min, the grain size decreased from 410.0 μm to 298.2 μm (a decrease of 27.3%), and the maximum corrosion depth after 8 days of 0.5% formic acid etching decreased from 74.9 μm to 58.3 μm (a decrease of 22.2%). The G2 / G1 ratio decreased to 2.7, indicating a more complete substructure and finer, more uniformly distributed recrystallized grains. This microstructure not only strengthens the pinning of grain boundaries by subgrain boundaries but also provides a dense matrix for the formation of a Cr2O3 / CrPO4 bilayer passivation film on the surface, thus improving high-temperature stability and corrosion resistance.

[0139] In Examples 13-16, the increased Ti content led to an increase in (G1+G2) / G3 from 3.0 to 124.0, a decrease in G2 / G1 from 45.9 to 2.4, an increase in ∑3+∑9 from 53.9% to 66.8%, and a decrease in ∑3 / ∑9 from 26.0 to 11.1, resulting in a refinement of the finished grain size from 5.9 μm to 3.5 μm. After heat treatment at 1000℃ for 30 min, the grain size decreased from 408.2 μm to 288.2 μm (a decrease of 29.4%), and the maximum corrosion depth after 8 days of 0.5% formic acid corrosion decreased from 75.1 μm to 59.0 μm (a decrease of 21.4%). The increased Ti content nearly eliminated G3 and made G2 absolutely dominant, resulting in a significant refinement of the finished grain size and thus improving the high-temperature stability of the alloy; at the same time, the formation of a Ti-based oxide film further improved the corrosion resistance.

[0140] In Examples 17-20, increasing the Zr content increased the (G1+G2) / G3 ratio from 3.5 to 165.7, decreased the G2 / G1 ratio from 39.9 to 2.1, increased the ∑3+∑9 ratio from 55.0% to 67.7%, and decreased the ∑3 / ∑9 ratio from 24.0 to 9.9, refining the finished grain size from 5.8 μm to 3.4 μm. After heat treatment at 1000℃ for 30 min, the grain size decreased from 392.7 μm to 272.0 μm (a decrease of 30.7%), and the maximum corrosion depth after 8 days of 0.5% formic acid corrosion decreased from 75.8 μm to 57.2 μm (a decrease of 24.5%). Zr further reduced the G2 / G1 and ∑3 / ∑9 ratios, and this optimized microstructure endowed the alloy with good high-temperature stability and corrosion resistance.

[0141] In Examples 21-29, increasing the Co content increased the (G1+G2) / G3 ratio from 4.0 to 199.0, decreased the G2 / G1 ratio from 35.4 to 1.7, increased the ∑3+∑9 ratio from 56.0% to 68.3%, and decreased the ∑3 / ∑9 ratio from 22.3 to 9.2, resulting in a refinement of the finished grain size from 5.5 μm to 3.2 μm. After heat treatment at 1000℃ for 30 min, the grain size decreased from 391.5 μm to 268.0 μm (a decrease of 31.5%), and the maximum corrosion depth after 8 days of 0.5% formic acid etching decreased from 74.5 μm to 54.0 μm (a decrease of 27.5%). The increase in Co content almost completely eliminates G3 (reducing from 20.0% to 0.5%), G2 becomes dominant, ∑3+∑9 increases to 68.3%, and ∑3 / ∑9 decreases to 9.2. This microstructure means that the deformation energy storage is almost completely released, the subgrain boundary pinning network is the densest, and the random grain boundary continuity is interrupted to the greatest extent. Both properties are the best of single elements.

[0142] Compared with Comparative Example 2 and Example 70, after adding element X, the volume percentage of G1 increased from 0.4% to 17.3%, the volume percentage of G2 increased from 23.6% to 79.1%, while the volume percentage of G3 decreased significantly from 76.0% to 3.6%. Furthermore, the proportions of ∑3 and ∑9 grain boundaries increased simultaneously, with the ∑3 / ∑9 ratio decreasing from 26.3 to 12.0. The finished grain size was refined from 5.5 μm to 4.5 μm. After heat treatment at 1000℃ for 30 min, the grain size decreased from 451.2 μm to 229.0 μm (a decrease of 49.2%), and the maximum corrosion depth after 8 days of 0.5% formic acid etching decreased from 92.3 μm to 48.7 μm (a decrease of 47.2%). This indicates that the introduction of element X changes the microstructure from deformation-dominated to recrystallization / substructure-dominated, greatly reducing the risk of abnormal grain growth driven by high strain energy at high temperatures. At the same time, the appropriate distribution of G1 and G2 enhances the subgrain boundary pinning network, and the significant reduction of ∑3 / ∑9 cuts off the continuity of random grain boundaries. High-temperature grain growth and the channels for corrosive media to penetrate along grain boundaries are simultaneously suppressed.

[0143] To investigate the effect of increasing X element types from mono- to hexa-ary on microstructure and properties, under conditions of similar total X content and a constant Co content of 0.1280%, the microstructure characteristics and target properties of mono-ary Co (Example 29), binary Fe-Co (Example 37), binary Cr-Co (Example 45), ternary Fe-Mn-Co (Example 57), quaternary Cr-Ti-Zr-Co (Example 63), pentagonal Mn-Cr-Ti-Zr-Co (Example 66), and hexa-ary Fe-Mn-Cr-Ti-Zr-Co (Example 69) were compared sequentially. As the X element type increased from mono- to hexa-ary, the core structure and grain boundary parameters remained highly stable. Among these, the G2 / G1 ratio remained within the optimized range of 1.7 to 5.1, the ∑3 / ∑9 ratio remained stable within a narrow range of 9.2 to 10.6, the (G1+G2) / G3 ratio remained between 65.7 and 199.0, and the proportion of special grain boundaries of ∑3+∑9 remained between 66.3% and 68.3%. Based on this, after treatment at 1000℃ for 30 min, the grain size decreased from 268.0 μm (unit Co) to 205.0 μm (hexa-membered Fe-Mn-Cr-Ti-Zr-Co), a reduction of 23.5%; the corrosion depth decreased from 54.0 μm to 43.5 μm, a reduction of 19.4%. This indicates that the introduction of multiple elements did not change the optimized microstructure established by Co, but rather added multiple benefits: different phosphides formed a multi-size, multi-distribution composite pinning network at the grain boundaries, creating multiple barriers to high-temperature grain boundary migration; and a dense film containing multiple complementary elements such as Co, Cr, Mn, and Fe was generated on the corroded surface. The phosphates / oxides of each element filled the defects, and the barrier ability of the film against the corrosive medium continued to increase with the increase of the number of elements.

[0144] Comparing the Co-free high-X content system (Examples 4, 8, 12, 16, 20, with a total X content of 0.2320~0.4788%) with the Co-containing low-X content system (Example 29, with a Co content of 0.1280%), the results show that the Co-containing system, with a lower X addition, is superior to the Co-free system in key microstructures. Its G3 is almost eliminated, the proportions of G1 and G2 are further optimized, and the proportion of ∑3+∑9 is further improved, resulting in better high-temperature grain size and etching depth.

[0145] Further comparison was made between the Co-containing binary low-X content systems (Example 37, Fe-Co system, total X 0.1295%; Example 45, Cr-Co system, total X 0.1296%) and the corresponding Co-free single-element high-X content systems (Example 4, Fe system, total X 0.4788%; Example 12, Cr system, total X 0.3429%). The high-temperature grain size of the Fe-Co system was 262.0 μm, and the corrosion depth was 55.0 μm, which were 24.1 μm and 19.5 μm smaller than those of the single Fe system (286.1 μm and 74.5 μm, respectively). The high-temperature grain size of the Cr-Co system was 261.0 μm, and the corrosion depth was 52.5 μm, which were both smaller than those of the single Cr system (298.2 μm and 58.3 μm, respectively). The above results demonstrate that the Co-containing binary system achieves superior high-temperature thermal stability and corrosion resistance with a total X content less than 40% of that in the Co-free monolithic system, fully showcasing the dual role of Co in enhancing high-temperature stability and optimizing the density of the corrosion film. This is because Co, even at a relatively low addition level, can effectively eliminate the G3 structure and improve the distribution of G1, G2, and special grain boundaries through its unique solute dragging effect and interfacial segregation behavior. Moreover, this structure-regulating effect cannot be replaced by simply increasing the Fe and Cr content.

[0146] Comparing Examples 69 and 70, when gradient annealing was used for the finished product, the proportion of G1 increased from 17.3% to 37.1%, the proportion of G2 decreased from 79.1% to 62.4%, and the proportion of G3 decreased from 3.6% to 0.5%; the proportion of ∑3 increased from 56.2% to 61.6%, and the proportion of ∑9 increased from 4.7% to 6.7%; the grain size of the finished product was refined from 4.5 μm to 2.8 μm. After heat treatment at 1000℃ for 30 min, the grain size decreased from 229.0 μm to 205.0 μm, a reduction of 10.5%; the maximum corrosion depth after 8 days of 0.5% formic acid etching decreased from 48.7 μm to 43.5 μm, a reduction of 10.7%. This demonstrates that the two-stage gradient annealing, through the synergistic effect of the recovery stage and the recrystallization stage, regulates the recovery driving force and recrystallization nucleation rate, thereby achieving directional optimization of grain structure distribution and special grain boundary distribution, and thus simultaneously improving the high-temperature thermal stability and corrosion resistance of the material.

[0147] According to the test results of Examples 21 to 29, the P / X ratio decreased from 115.0 to 1.8. As the P / X ratio decreased, the proportion of G1 increased from 2.2% to 36.9%, the proportion of G2 increased from 77.8% to 91.6% and then fell back to 62.6%, and the proportion of G3 decreased from 20.0% to 0.5%, reflecting that a lower P / X ratio is beneficial for eliminating residual deformation structures and regulating the recrystallization driving force. The ratio of (G1+G2) / G3 increased from 4.0 to 199.0, and the ratio of G2 / G1 decreased from 35.4 to 1.7, indicating that the microstructure composition shifted from being dominated by substructures to being dominated by recrystallized states. Regarding grain boundary characteristics, the proportion of ∑3 increased from 53.6% to 61.6%, the proportion of ∑9 increased from 2.4% to 6.7%, and the total of ∑3+∑9 increased from 56.0% to 68.3%. The ∑3 / ∑9 ratio decreased from 22.3 to 9.2, with the increase in ∑9 being more significant. This indicates that a lower P / X ratio promotes multiple special grain boundary reactions, forming a higher density of special grain boundary networks and reducing the continuity of random grain boundaries. The finished grain size was refined from 5.5 μm to 3.2 μm; after heat treatment at 1000℃ for 30 min, the grain size decreased from 391.5 μm to 268.0 μm; after etching in 0.5% formic acid solution for 8 days, the maximum corrosion depth decreased from 74.5 μm to 54.0 μm (see Tables 2-3 and 4). Figure 1 When the P / X ratio is in the range of 1 to 10 (Examples 26-29), the alloy recrystallization is more complete, the distribution of grain boundary features is optimized, and the ratio of ∑3 to ∑9 is at a reasonable level. This provides sufficient grain boundary pinning sites while avoiding the introduction of too many high-energy interfaces due to an excessively high ∑9 ratio, which would affect the high-temperature grain size stability and corrosion resistance. Therefore, the high-temperature grain size stability and corrosion resistance of the alloy are better matched. When the P / X ratio is in the range of 10 to 25, the degree of recrystallization of the alloy is relatively low, the proportion of residual deformed structure increases, the ∑9 ratio decreases, the ∑3 / ∑9 ratio increases, the density of special grain boundary networks is relatively low, the connectivity of random grain boundaries is enhanced, and the grain growth resistance and corrosion resistance decrease accordingly.

[0148] The (G1+G2) / G3, G2 / G1, ∑3+∑9, ∑3 / ∑9, and finished grain sizes of Example 16 (Ti, 0.2709%) and Example 12 (Cr, 0.3429%) are highly similar. Based on this microstructure, the high-temperature grain size of the Cr system (298.2 μm) is 10 μm larger than that of the Ti system (288.2 μm), and the maximum corrosion depth (58.3 μm) is 0.7 μm smaller than that of the Ti system (59.0 μm). The above comparison shows that, under similar microstructure conditions, the high-temperature grain size stability of phosphides / oxides in the Cr system is weaker than that of Ti-based compounds. That is, the thermal stability of the element type and the chemical stability of the film layer affect the final service performance. Therefore, the optimization of the microstructure and the selection of element types need to be considered together in alloy design.

[0149] This invention selectively introduces Fe, Mn, Co, Ti, Cr, and Zr elements, synergistically controlling their addition ratio with P, and combining this with optimized production processes to ultimately produce copper alloy tubes that possess excellent thermal stability, good corrosion resistance, and machinability, while maintaining controllable production costs. Based on these comprehensive performance advantages, this tube can be widely used in heat exchange, new energy, and high-power computers, and is particularly suitable for products such as heat exchange tubes, condensers, and evaporators that require brazing and operate in corrosive environments. This invention allows for flexible adjustment of the types and proportions of added elements according to different operating conditions, achieving effective adaptation between material properties and application scenarios.

[0150] It should be noted that the above specific embodiments are merely preferred embodiments of the present invention, used to exemplify the technical solution and its effects, and are not intended to exhaustively list all possible embodiments of the present invention. Those skilled in the art can test other binary or multi-element combinations not listed in the embodiments as needed.

[0151] Table 1. Alloy composition (mass percentage %) of the examples and comparative examples.

[0152] Table 2 Microstructure parameters of alloys in the examples and comparative examples

[0153] Table 3. Performance comparison of the alloys in the examples and comparative examples.

Claims

1. A corrosion-resistant and high-temperature resistant copper alloy pipe, composed of P, X, copper, and unavoidable impurities, wherein X is selected from at least one of Fe, Mn, Cr, Ti, Zr, and Co, characterized in that, By mass percentage, 0.12%≤P≤0.36%, 0.0012%≤Fe≤0.50%, 0.0015%≤Mn≤0.45%, 0.0012%≤Cr≤0.40%, 0.0015%≤Ti≤0.30%, 0.0013%≤Zr≤0.27%, 0.0015%≤Co≤0.15%, the total content of X satisfies 0.0012%≤X≤0.50%, the total amount of unavoidable impurities is ≤0.01%, and copper is the balance; The copper alloy tube comprises recrystallized grain structure G1, substructured grain structure G2, and deformed grain structure G3, with the following volume percentages: 0.50%≤G1≤40.0%, 55.0%≤G2≤95.0%, and 0.30%≤G3≤43.0%. G1 consists of grains with a grain orientation dispersion ≤1°, G2 consists of grains with a grain orientation dispersion ≤5° and 1° < grain orientation dispersion, and G3 consists of grains with a grain orientation dispersion >5°.

2. The corrosion-resistant and high-temperature-resistant copper alloy pipe according to claim 1, characterized in that, The average grain size of the copper alloy tube is 1.0~20.0μm.

3. The corrosion-resistant and high-temperature-resistant copper alloy pipe according to claim 1, characterized in that, In the copper alloy tube, the proportion of ∑3 grain boundaries satisfies 40.0%≤∑3≤70.0%, and the proportion of ∑9 grain boundaries satisfies 1.0%≤∑9≤8.0%.

4. The corrosion-resistant and high-temperature-resistant copper alloy pipe according to claim 1, characterized in that, The mass ratio of P to X satisfies: 0.3 ≤ P / X ≤ 180.

0.

5. The corrosion-resistant and high-temperature-resistant copper alloy pipe according to claim 1, characterized in that, The mass ratio of P to X satisfies: 0.4 ≤ P / X ≤ 177.

0.

6. The corrosion-resistant and high-temperature-resistant copper alloy pipe according to claim 1, characterized in that, The X contains Co, where 0.0020% ≤ Co ≤ 0.15%.

7. The corrosion-resistant and high-temperature-resistant copper alloy pipe according to claim 1, characterized in that, The X contains Fe and Mn, with 0.0015%≤Fe≤0.06% and 0.0018%≤Mn≤0.13%.

8. The corrosion-resistant and high-temperature-resistant copper alloy pipe according to claim 1, characterized in that, The X contains 0.01%≤Co≤0.13% and satisfies 0.0015%≤[Fe, Mn, Cr, Ti, Zr]≤0.06%, where [Fe, Mn, Cr, Ti, Zr] represents at least one selected from Fe, Mn, Cr, Zr, and Ti.

9. The corrosion-resistant and high-temperature-resistant copper alloy pipe according to claim 8, characterized in that, In the [Fe, Mn, Cr, Ti, Zr], 0.0015%≤Fe≤0.05%, 0.0015%≤Mn≤0.025%, 0.0015%≤Cr≤0.026%, 0.0015%≤Ti≤0.032%, and 0.0015%≤Zr≤0.027%.

10. The corrosion-resistant and high-temperature-resistant copper alloy tube according to claim 9, characterized in that, The X contains Co and Zr, with 0.05% ≤ Co ≤ 0.13% and 0.0015% ≤ Zr ≤ 0.027%.

11. The corrosion-resistant and high-temperature-resistant copper alloy pipe according to claim 1, characterized in that, The X comprises Co, 0.01% ≤ Co ≤ 0.13%, and elements selected from group (a) or group (b) below: (a) 0.0015%≤Fe≤0.012%, 0.0015%≤Mn≤0.022%; (b) 0.0015% ≤ Ti ≤ 0.012%, 0.0015% ≤ Zr ≤ 0.011%.

12. The corrosion-resistant and high-temperature-resistant copper alloy pipe according to claim 1, characterized in that, The X contains Co, Cr, Ti and Zr, with 0.08%≤Co≤0.13%, 0.0015%≤Cr≤0.006%, 0.0015%≤Ti≤0.011%, and 0.0015%≤Zr≤0.007%.

13. The corrosion-resistant and high-temperature-resistant copper alloy pipe according to claim 1, characterized in that, The X contains Co, Mn, Cr, Ti and Zr, with 0.09%≤Co≤0.13%, 0.0015%≤Mn≤0.011%, 0.0015%≤Cr≤0.005%, 0.0015%≤Ti≤0.009%, and 0.0015%≤Zr≤0.0065%.

14. The corrosion-resistant and high-temperature-resistant copper alloy pipe according to claim 1, characterized in that, The X comprises Co, Fe, Mn, Cr, Ti, and Zr, with 0.10%≤Co≤0.13%, 0.0015%≤Fe≤0.008%, 0.0015%≤Mn≤0.0065%, 0.0015%≤Cr≤0.005%, 0.0015%≤Ti≤0.0068%, and 0.0015%≤Zr≤0.004%.

15. The corrosion-resistant and high-temperature-resistant copper alloy pipe according to claim 1, characterized in that, The volume ratio of G1, G2 and G3 satisfies: (G1+G2) / G3≥1.5; And / or, the volume ratio of G2 to G1 satisfies: 1.5≤G2 / G1≤65.

0.

16. The corrosion-resistant and high-temperature resistant copper alloy pipe according to claim 1, characterized in that, In the copper alloy tube, the sum of the proportions of ∑3 grain boundaries and ∑9 grain boundaries satisfies: 46.0% ≤ ∑3 + ∑9 ≤ 70.0%; And / or, the ratio of the proportions of ∑3 grain boundaries to ∑9 grain boundaries satisfies: 8.0 ≤ ∑3 / ∑9 ≤ 35.

0.

17. The corrosion-resistant and high-temperature-resistant copper alloy pipe according to claim 1, characterized in that, In the copper alloy tube, the ratio of ∑3 grain boundaries to ∑9 grain boundaries satisfies: 8.0≤∑3 / ∑9≤15.

0.

18. The corrosion-resistant and high-temperature-resistant copper alloy tube according to any one of claims 1-17, characterized in that, The average grain size of the copper alloy tube after heat treatment at 1000℃ for 30 min is 195~420μm, and the maximum corrosion depth after alternating hot and cold corrosion in a 0.5% volume fraction formic acid aqueous solution atmosphere for 8 days is 40~78μm.

19. The method for preparing the corrosion-resistant and high-temperature resistant copper alloy tube according to any one of claims 1-18, characterized in that, The process includes the following steps: batching and smelting → horizontal continuous casting → planetary rolling → stretching → finished product annealing; In the planetary rolling process, the deformation is 90-92%, and after rolling, the resulting tube blank is cooled to below 100°C within 0.5-1.5 seconds. In the finished product annealing process, gradient annealing is adopted. In an inert atmosphere, the pipe is held at 260~290℃ for 40~80 minutes and then held at 570~595℃ for 30~50 minutes.

20. The preparation method according to claim 19, characterized in that, In the batching and smelting process, the materials are fed in a reverse order according to their reactivity: S1: Add copper raw materials, and when the melt temperature reaches 1150~1230℃, add Cu-P master alloy according to the ratio, and hold for 10~15 minutes. S2: Reduce the melt temperature to 1100~1140℃, and selectively add intermediate alloys containing Co, Fe, Mn, Cr, Ti and Zr in sequence according to the formula. After each intermediate alloy is added, hold the temperature for 5~10 minutes.

21. The preparation method according to claim 19, characterized in that, In the stretching process, there are at least 4 stretching passes, with the elongation coefficient of the first 2 passes being 1.35~1.65; and the elongation coefficient of subsequent passes being 1.15~1.30, while controlling the pipe temperature to not exceed 50℃.

22. The preparation method according to claim 19, characterized in that, The process, following the stretching step and preceding the finished product annealing step, also includes online annealing and internal thread forming. In the online annealing step, the online annealing temperature is 530~850℃, the feed rate is 330~430m / min, and the online annealing temperature T and feed rate v are controlled to satisfy the following relationship: T × v = 2.0×10⁻⁶ 5 ~ 2.8×10 5 ℃·m / min.

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