KR101476753B1 - Method for producing an iron-chromium alloy - Google Patents

Abstract

The present invention relates to a process for thermomechanically treating a semi-finished product comprising a Laves phase and / or an Fe-containing particle and / or a Cr-containing particle and / or a Si-containing particle and / or an iron- Wherein the alloy is subjected to solution heat treatment at a temperature above the solidification heat treatment temperature in the first step and is then subjected to a solution heat treatment at a temperature above the solidification heat treatment temperature, Or in air or in water. In the second step, a mechanical forming of the semi-finished product is performed in the range of 0.05 to 99%, and the parts produced from the semi-finished product molded in the subsequent step are heated at a heating rate of 0.1 캜 / min to 1000 캜 / min by being heated up to the application temperature of 550 ℃ to 1000 ℃, Laves phase Fe ₂ (M, Si) or Fe ₇ (M, Si) ₆ and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles and / Or the carbide is deposited in a finely distributed manner as desired.

Description

FIELD OF THE INVENTION [0001] The present invention relates to a method for producing an iron-chromium alloy,

The present invention relates to a ferritic iron-chromium alloy made of molten metallurgy.

From DE 100 25 108 A1 it is known that from the Y, Ce, Zr, Hf and Al groups up to 2% by weight of one or more oxygen affinitive elements, at high temperature with chromium oxide in the spinel phase of the MCr ₂ O ₄ type 2% by weight or less of M element from the Mn, Ni and Co groups forming the Ti, Hf, Sr, Ca and Zr groups which increase the electrical conductivity of the Cr-based oxide, Lt; RTI ID = 0.0 > of iron < / RTI > The chromium content of this German publication is in the range of 12% to 28%. The field of application for these high temperature materials is the bipolar plates of high temperature fuel cells.

EP 1 298 228 A1 relates to a steel for a high temperature fuel cell wherein the steel contains less than 0.2% C, less than 1% Si, less than 1% Mn, less than 2% Ni, 15-30% Cr And a composition containing 1% or less of Al, 0.5% or less of Y, 0.2% or less of SE, 1% or less of Zr and the balance iron and impurities according to manufacturing conditions.

The common point of the two alloys is that they have low heat resistance and insufficient creep resistance at temperatures from 700 ° C. Above all, however, the alloy has excellent oxidation resistance and corrosion resistance in the region from above 700 ° C to about 900 ° C.

A ferritic steel having a high creep strength is known from DE 10 2006 007 598 A1. This steel is a metal alloy of the type Fe ₂ (M, Si) or Fe ₇ (M, Si) ₆ containing at least one metal alloy element M which can be formed by niobium, molybdenum, tungsten or tantalum elements intermetal phase. The steel according to the German publication is preferably used as a bipolar plate of a fuel cell stack.

EP 1 536 031 A1 discloses metal materials for fuel cells. The metal material may be selected from the group consisting of 0.2% or less of C, 0.02 to 1% of Si, 2% or less of Mn, 10 to 40% of Cr, 0.03 to 5% of Mo, 0.1 to 3% of Nb, Y, La, Ce, Pr, Nd, Pm, Sn, Zr and Hf groups, the balance iron and unavoidable impurities, wherein the composition satisfies the equation 0.1 Mo / Nb 30 .

EP 1 882 756 A1 discloses ferrite chromium steel, which is particularly useful in fuel cells. The chrome steel is characterized by comprising at most 0.1% C, 0.1-1% Si, 0.6% Mn, 15-25% Cr, 2% Ni, 0.5-2% Mo, 0.2-1.5% Nb, Ti, 0.5% Zr, 0.5% Zr, 0.3% SE, 0.1% Al, 0.07% N, balance Fe and impurities according to melting conditions, wherein the Zr + Ti content is 0.2 %.

All of these alloys have improved heat resistance and elevated creep resistance at temperatures from 700 DEG C compared to DE 100 25 108 A1 and EP 1 298 228 A2, and more precisely the dislocation motions of the material and the corresponding plastic deformation The heat resistance and the creep resistance increase through the formation of precipitation which prevents the formation of the precipitate. The precipitate is, for example, in the case of DE 10 2006 007 598 A1, composed of an intermetallic compound having a composition of Laveth phase, that is, Fe ₂ (M, Si) or Fe ₇ (M, Si) ₆ , M is niobium, molybdenum , Tungsten, or tantalum. In this case, a volume percentage of 1 to 8%, preferably 2.5 to 5%, is achieved. However, the precipitate may be another precipitate such as Fe containing particles and / or C containing particles and / or Si containing particles as disclosed in EP 1 536 031 A1, or may be a carbide containing Nb, W, Mo have. The common point of all of these particles is that it is difficult to deform the material.

It is known from the above-mentioned conventional art that the addition of a small amount of reactive elements such as Y, Zr, Ti, Hf, Ce, La and the like can very positively affect the oxidation resistance of the Fe-Cr alloy.

The alloys listed in DE 10 2006 007 598 A1, EP 1 536 031 A1 and EP 1 882 756 A1 are optimized for application as interconnector plates for high temperature fuel cells. In other words, the alloy exhibits a coefficient of expansion that is suitably adjusted to the ceramic component anode and electrolyte through the use of a ferrite alloy containing 10 to 40% chromium.

An additional requirement for inter-connector steels for high temperature fuel cells is the excellent corrosion resistance, good conductivity of the oxide layer, and low chromium evaporation in addition to the previously mentioned creep strength.

The requirements for reformers and heat exchangers for high temperature fuel cells are as high as possible creep strength, very good corrosion resistance, and low chromium evaporation. The oxide should not be conductive in the case of the component.

For example, the requirements for exhaust gas conduits in internal combustion engines, or parts for steam boilers, superheaters, turbines and other parts of power plants are as high as possible creep strength and very good corrosion resistance. In this case, the chromium evaporation does not cause a poisoning symptom as in a fuel cell, and the protective oxide should not be conductive in the case of the parts.

In DE 10 2006 007 598 A1, for example, excellent corrosion resistance is achieved through the formation of a chromium oxide coating. Spinel containing Mn, Ni, Co or Cu is further formed on the chromium oxide coating layer, thereby forming chromium oxide or chromium oxyhydroxide with lower volatility that poisons the cathode. Further, through the stable bonding of Si in Lavess phase Fe ₂ (M, Si) or Fe ₇ (M, Si) ₆ , a nonconductive underlayer made of silicon oxide is formed below the chromium oxide coating layer It does not. Corrosion resistance is further enhanced through the fact that the Al content is kept low and thus the increase in corrosion is prevented by the internal oxidation of aluminum. In addition, a small amount of Ti addition provides hardening of the surface, thereby preventing the bulging of the oxide layer and the inclusion of the metal regions in the oxide layer, thereby increasing the oxidation. In addition, the addition of oxygen affinity elements such as La, Ce, Y, Zr and the like further increases corrosion resistance.

In the market, it is required to increase the application temperature and the heat resistance and creep strength at an elongation of 18% or more and the higher use temperature of the alloy in order to prevent brittle fracture in a state where the oxidation resistance or corrosion resistance is maintained at least at the same level, More precisely, there is an increasing demand for products that must maintain acceptable deformability, measured as plastic deformation in tensile tests at room temperature and elongation in excess of 13%.

The following analysis method is also used.

In the creep test, the sample is stretched at a constant static tensile force under constant temperature conditions. The tensile force is indicated as the initial tensile stress in relation to the initial sample cross-section for comparison. In the creep test, the simplest case is to measure the time t _B (breaking time) until the sample is destroyed. The test can thus be carried out without measuring the elongation of the sample during the test. The fracture elongation is then measured after the end of the test. For example, the sample is loaded into a creep tester at room temperature and heated to the desired temperature under conditions where no load is applied by tensile force. After reaching the test temperature, the sample remains unchanged for 1 hour without load application for temperature equilibrium. The sample is then subjected to a tensile force and the test time begins.

The fracture time can be used as a measure of creep strength. The longer the fracture time at a given temperature and initial tensile stress, the higher the creep strength of the material. The breakdown time and creep strength decrease as the temperature rises and the initial tensile stress increases (see, for example, "Buergel's High Temperature Material Handbook," p. 100).

The deformability is measured by a tensile test according to DIN 50145 at room temperature. In this case it is measured elastic limit R _{p0 .2,} a tensile strength R _M and elongation to fracture. The elongation (A) is measured from the elongation length of the initial measured length L ₀ in the fractured sample, and the formula is as follows.

In the above equation, L _U is the measured length after fracture.

Depending on the length of each measurement, the fracture elongation is given by the subscript as shown below.

A ₅ , measurement length L ₀ = 5 · d ₀ or L ₀ = 5.65 · S ₀ ,

A ₁₀ , the measurement length L ₀ = 10 · d ₀ or L ₀ = 11.3 · S ₀ ,

Or A _{L = 100} if, for example, the randomly chosen measurement length L = 100 mm.

(d ₀ is the initial diameter and S ₀ is the initial cross-sectional area of the flat sample.)

The value of elongation (A) in a tensile test at room temperature can be used as a measure of deformability.

The Laveth phase (s) or the Fe-containing particles and / or Cr-containing particles and / or the Si-containing particles and / or carbides may be removed by etching using a V2A-corrosive (V2A-Beize) or electrolytic etching using oxalic acid Can be visualized. Grain or grain boundaries are additionally etched at the time of etching with V2A caustic. In this case, only particles of about 0.5 mu m in size are observed when observed under an optical microscope. Smaller particles can not be identified, but of course they can exist. Therefore, metallography is used only as an auxiliary means for explanation, and in a more practical sense, the effectiveness of the method is evaluated by the breaking time and the creep strength.

In Ralf Buergel's Handbuch der Hochtemperaturwerkstofftechnik, published in December 2006 by Viehweg Publishers, and in the third edition ("Verger"), pages 196-199 and Table 3.7, Possible measures to increase the creep strength of the metallic material "are described.

Action methods,

- "high melting point, face-centered cubic material"

- "high modulus of elasticity",

- "Materials with low stacking fault energies

Is not used for the improvement of the parameters raised above because the action requires a change in the material type and this is not possible here nor for the purposes of this application.

Action methods,

"Solid solution hardening ", "

"Particle hardening ", "

- "high particle volume percent",

- "The particles with a low diffusion coefficient of the corresponding alloying element"

DE 10 2006 007 598 A1 and / or EP 1 536 031 A1 and / or EP 1 882 756 A1.

Action methods,

- "particles with low solubility in the matrix",

- "Coherent particles with low interfacial enthalpy for the matrix"

It does not apply to the precipitates observed.

Similarly,

- "carbides or borides as grain boundary precipitates, oxides and sulfides prevention ", "

- "Addition of precisely controlled quantities of positive working elements such as B, C, Zr and Ce"

- "higher purity of the alloy",

- "getter element (eg S) addition",

- "High corrosion resistance"

Have already been described in DE 10 2006 007 598 A1 and / or EP 1 536 031 A1 and / or EP 1 882 756 A1.

Solution

- "dendrite arm spacing" for castings,

- "Grain structure oriented toward the main load direction",

- "single crystal",

- "Low densities for parts that rotate under magnetic weight"

It does not apply to this alloy type or manufacturing method or use.

For the purpose of improving the creep strength of precipitation hardened iron-chromium alloy,

1) "coarse grain texture ", "

2) "interlocking of grain boundaries by precipitates"

3) "Optimized heat treatment (optimum grain size control, elimination of segregation in the casting structure, optionally controlled grain boundary roughness)",

4) "Cold deformation prevention" is considered.

It is an object of the present invention to provide a method for producing a component made of precipitation hardened iron-chromium alloy, which has a higher heat resistance and creep strength of the precipitation hardened ferrite alloy, while maintaining allowable deformability at room temperature, And a method for producing the same.

It is a further object of the present invention to provide a component / semi-finished product which is thermomechanically treated and made of iron-chromium alloy and which can be used to achieve high heat resistance and creep strength, respectively, while maintaining acceptable deformability at room temperature, Semi-finished products.

A final object of the present invention is to make available the parts / semi-finished products manufactured as described above for specific technical applications in a temperature range exceeding 550 ° C.

The above object is also achieved by a process for producing a semi-finished product made of an iron-chromium alloy for depositing a Laves phase and / or an Fe-containing particle and / or a Cr-containing particle and / A method of manufacturing a component made of the alloy by thermomechanical treatment,

In the first step, the alloy is subjected to solution heat treatment at a temperature above the solidification heat treatment temperature and then cooled in a protective gas or air that is stationary, in a protective gas or air that is traveling (blowing) ,

In the second stage, the semi-finished product is subjected to mechanical forming in the range of 0.05 to 99%, and in the subsequent stage, the semi-finished product is heated at a heating rate of 0.1 ° C / min to 1000 ° C / min by applying load to the application temperature of 550 ℃ to 1000 ℃, Laves phase Fe ₂ (M, Si) or Fe ₇ (M, Si) ₆ and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles And / or the above-described parts manufacturing method in which the carbide is deposited in a targeted and finely distributed manner.

On the other hand, the above object can be achieved by a process for heat-treating a semi-finished product made of an iron-chromium alloy for precipitating a Lavess phase and / or an Fe-containing particle and / or a Cr-containing particle and / A method of manufacturing a component comprising:

In the first step, the alloy is solid solution heat treated at a temperature above the solidification heat treatment temperature and then cooled in a protective gas or air which is stationary, in a protective gas or air moving (blowing)

In the second step the mechanical molding of the semi-finished product is carried out in the range of 0.05 to 99%, and in the subsequent step the molded semi-finished product is heated in the temperature range of 550 to 1060 캜 under protective gas or air for a period of time from t _min to t _max Cooled in a protective gas or air that is heat-treated and then stopped, or in protective gas or air that is moving (blowing), or in water, or cooled to less than 800 ° C in a furnace for heat treatment, The Fe-containing particles and / or the Cr-containing particles and / or the Si-containing particles and / or the carbides are dispersed in a targeted and finely distributed manner, such as Fe ₂ (M, Si) or Fe ₇ (M, Si) ₆ and / And the target part is achieved by a method of manufacturing a part which is completed before or after the heat treatment.

Where t _min and t _max are calculated according to the following formula:

및

이고, T_a = T + 273.15이다.

And

And T _a = T + 273.15.

The times t _min and t _max are in minutes and the unit of heat treatment temperature (T) is in ° C.

Advantageous developments of the method according to the invention can be found in the dependent claims according to the corresponding method.

In the first step, the following temperature ranges and times are provided for the heat treatment of the solid solution.

≫ 1050 < 0 > C (> 1050 < 0 &

-> 1060 ° C (> 1060 ° C) over a period of more than 1 minute.

The resultant changes in material properties are explained in more detail in the course of further explanation.

The above object is also achieved by a metal part or a semi-finished product,

(By weight) of the following chemical composition,

Cr 12 - 30%

Mn 0.001 - 2.5%

0.1 to 2% of Nb,

W 0.1 - 5%,

0.05 - 1% of Si,

C 0.002 - 0.1%

N 0.002 - 0.1%

S Up to 0.01%

Fe rest, and

Unavoidable molten impurities,

After the processing heat treatment, the Laveth phase (s) exhibit a deformed tissue structure deposited in a finely distributed form in the tissue dislocation of the tissue structure, and at a stretching rate of 750 ° C and 18% or more, for example, 35 MPa In the creep test used is achieved by a metal part or semi-finished product which is set within the tissue structure by a breaking time which exceeds 1.5 times the breaking time of the coarse-grained fully recrystallized tissue.

Similar results are achieved in creep tests with different stresses and temperatures, where the temperature for the creep test is preferably in the range of 500 to 1000 占 폚.

In the following, the above-described measures 1 to 4 are considered.

Surprisingly, preforming followed by a suitably adjusted annealing treatment as compared to measure 4 "cold deformation" results in a creep test in which the fracture time for coarse grain structure (action method 1) exceeds 1.5 times, It is possible to achieve an extension of the breakdown time of the sample, preferably exceeding 3 times.

In this case, in the case of the third step (precipitation of Lavess phase (s)), the molded semi-finished products or parts made of this semi-finished product, as the case may be, is heated to a heat treatment temperature of ℃ the combination is heat treated for a period of time t _min to t _max at the heat treatment temperature under a protective gas or air in a subsequent stage, followed by a stationary protective gas or in air, movement (blown) protective gas to Or cooled in air or in water, or in a furnace for a heat treatment to 800 ° C or lower, and then, if desired, the desired part is produced. T _min and t _max are calculated according to the following formula.

및

이고, T_a = T + 273.15이다.

And

And T _a = T + 273.15.

In addition, for the third stage (precipitation on Laveth), the semi-finished or semi-finished parts are heat treated under protective gas or air in the temperature range from 550 to 1060 ° C for a period of from t _min to t _max , Cooled to below 800 ° C in a furnace, in a gas or air, in a protective gas or air moving (blowing), or in water, or for heat treatment, and then the desired component is manufactured, The next completed part can be heated to an application temperature of 550 ° C to 1000 ° C through heating at 0.1 ° C / min to 1000 ° C / min. Herein, t _min And t _max are calculated according to the following formula:

및

이고, T_a = T + 273.15이다.

And

And T _a = T + 273.15.

According to another aspect of the present invention,

A semi-finished product consisting of an alloy with the following chemical composition (by weight%) is processed by a processing heat treatment:

Cr 12 to 30%

Mn 0.001 to 2.5%

0.1 to 2% Nb,

W 0.1 to 5%

0.05 to 1% of Si,

C 0.002 to 0.03%

N 0.002 to 0.3%

S Up to 0.01%

Fe rest, and

Unavoidable molten impurities.

By means of the process according to the invention, semi-finished products are produced in the form of plates, strips, rods, forgings, tubes or wires, and the parts are finished in various forms required for each application.

Particularly preferably, it is subjected to a solid solution heat treatment at a temperature equal to or higher than the solidification heat treatment temperature, preferably at a temperature of 1050 占 폚 or more for each of more than 6 minutes, or at a temperature of 1060 占 폚 or more for 1 minute or more, or in the air, moving from (blowing are) protective gas or air, or after cooling in water, the LA in the alloy bath in an initial state prior to the forming of the said semi-finished Fe ₂ (M, Si) or Fe ₇ ( M, Si) ₆ and / or Fe-containing particles and / or Cr-containing particles or Si-containing particles and / or carbides.

Molding of semi-finished products can be accomplished through hot forming. Alternatively, however, the molding may also be carried out through cold forming.

In the first case, the semi-finished product is hot formed to a starting temperature of more than 1070 ° C, and the final mechanical strain of 0.05 to 95% is carried out at 1000 ° C to 500 ° C, preferably the last 0.5 to 90% Lt; / RTI >

In the second case, the degree of cold forming of the semi-finished product is 0.05 to 99%, preferably 0.05 to 95% or 0.05 to 90%.

According to another embodiment of the present invention, the mechanical molding of the semi-finished product is 20 to 99%, and then the molded semi-finished product is heated at a temperature ranging from 950 캜 to 1060 캜 under protective gas or air for a period of t _min to t _max It is proposed that the object is cooled in a protective gas or air that is heat-treated and then stopped, in a moving protective gas or air, or in air, or in water, after which the target component is completed, wherein t _min and t _max are It is calculated according to the following formula:

및

이고, T_a = T + 273.15이며, t_min 및 t_max의 단위는 분이고, 상기 열처리 온도(T)의 단위는 ℃이다. 앞서 명시한 합금이 고체 산화물 연료 전지를 위한 인터커넥터로서 이용되는 경우, 알루미늄 함량은 0.001 - 0.5%가 바람직하다.

And

T _a = T + 273.15, the units of t _min and t _max are minutes, and the unit of the heat treatment temperature (T) is ° C. When the above-described alloy is used as an interconnector for a solid oxide fuel cell, the aluminum content is preferably 0.001 - 0.5%.

In another application, for example a reformer or a heat exchanger for a fuel cell in which a conductive oxide layer is not needed, the aluminum content is preferably 2 to 6%, which is better than chromium-manganese spinel Again a closed aluminum oxide layer with a lower growth rate and additionally lower chromium oxide evaporation can again be formed.

Both of these variations can be considered in the field of applications where there is no need for conductive oxides and no special requirement for chromium evaporation. A particular consideration here is that as the aluminum content increases, the processability and weldability of the alloy deteriorate, resulting in higher costs. Therefore, when the oxide layer is composed of chromium oxide and chromium-manganese spinel, it is possible to ensure sufficient oxidation resistance by using 0.001-0.5% aluminum. When a higher oxidation resistance is required, for example, as guaranteed through the formation of an aluminum oxide layer, the aluminum content is preferably 2.0 to 6.0%. The two alloy variants can be used, for example, as exhaust gas conduits in internal combustion engines, or as components for steam boilers, superheaters, turbines and other parts of power plants.

The range of preferred aluminum contents is in the range of 2.5% to 5.0%, which is characterized by particularly good processability.

The following alloys may additionally incorporate the following elements individually or in combination with each other.

0.02 to 0.3% of La,

0.01 to 0.5% of Ti,

0.0001 to 0.07% Mg,

Ca 0.0001 to 0.07%

P 0.002 to 0.03%

0.01 to 3% Ni / Co / Cu,

B 0.005% or less.

The content of elements which can additionally be incorporated in the alloy may be set to 0.0001 to 0.05% Mg, 0.0001 to 0.03% Ca, 0.002 to 0.03% P,

In addition, the alloy may contain at least one of Ce, La, Pr, Ne, Sc, Y, Zr or Hf in an amount of 0.02 - 0.3% (by wt.

If necessary, the alloy may contain 0.02 - 0.2% of at least one of Ce, Pr, Ne, Sc, Y, Zr or Hf elements (by weight%).

In order to achieve the desired effect, the Nb content is 0.3 to 1.0% and the Si content is 0.15 to 0.5%.

If necessary, the tungsten element may be completely or partially replaced by at least one of Mo or Ta elements.

In addition, if necessary, the alloy may contain up to 0.2% V and / or up to 0.005% S. In this case, the oxygen content should not be greater than 0.01%.

Further, if necessary, the alloy may contain a maximum of 0.003% of boron.

In addition, the alloy contains Zn, Sn, Pb, Se, Te, Bi, and Sb elements at a maximum of 0.01%.

On the other hand, the part / semi-finished product which is composed of the above-mentioned alloy composition and which is manufactured according to the method of the present invention is preferably used as an interconnector in a fuel cell or as a material of a part such as a reformer or heat exchanger in an auxiliary device of the fuel cell As shown in FIG.

Alternatively, the part / semi-finished product or alloy itself manufactured in accordance with the process of the present invention may be used as a component in the exhaust gas conduit of an internal combustion engine, or as a component for steam boilers, superheaters, turbines and other parts in power plants or chemical process industries There is also the possibility of being used as a component.

Due to the processing heat treatment in the alloys produced by the molten metallurgical process through the process according to the invention, the Laves phases can be precipitated in a manner finely distributed as desired in the potential part of the structure.

The details and advantages of the invention are explained in more detail in the following examples.

The steps of the method according to the invention will now be discussed in more detail.

The first step of the iron-chromium alloy for precipitating the Lavess phase and / or the Fe-containing particles and / or the Cr-containing particles and / or the Si-containing particles and / or the carbides in the processing heat treatment is a Laves phase and / Annealing above the solidification heat treatment temperature must be carried out in order to allow the Cr-containing particles and / or the Si-containing particles and / or carbides to dissolve and be used for precipitation for subsequent processing heat treatment. The heat treatment temperature for the solidification treatment is determined depending on the alloy, but is preferably more than 1050 DEG C for a period of more than 6 minutes, or more than 1060 DEG C for a period of more than 1 minute, Cooling) in protective gas or air, or in water. Exact temperature control above the solidification heat treatment temperature is not determined with respect to the characteristics. The annealing can be done in air or under protective gas. However, the annealing should be carried out below the melting temperature, preferably below 1350 ° C. Also, the annealing time should preferably be less than 24 hours for cost reasons, but may be longer depending on the alloy. Following the solidification heat treatment, cooling is carried out in stationary protective gas or air, in moving (blowing) protective gas or air, or in water, at which time only a few Laves phases are formed.

In addition, it should be noted that, in particular for parts of thick walls, all parts of the part must reach the minimum annealing time required at the preset temperature conditions. This should be taken into consideration when determining the starting point of the annealing time.

In the second step, an increased dislocation density should be introduced into the material. Increased dislocation densities have deformed or restored tissue where the dislocations are arranged at small-angle grain boundaries.

Thus, the second step is to provide a uniform distribution of the Laveth phase and / or the Fe-containing particles and / or the Cr-containing particles and / or the Si-containing particles and / or the carbides in the subsequent annealing treatment, to be less to that sex-type.

Such modifications may be cold forming, but may also be hot forming, and in hot forming, the structure must not be completely recrystallized at the time of rolling. This is accomplished by limiting the strain zone for the final molding and the temperature at which the final molding takes place. Upon deformation of more than 1000 DEG C, the material tends to be already recrystallized or recovered during the shaping, so that the shaping should preferably take place at a temperature below 1000 deg. There is brittleness in the temperature range of less than 500 DEG C and 475 DEG C in the case of ferrite. In this case the elongation is lower and the molding resistance is increased, which causes somewhat undesirable molding and reduces the economics.

Precipitates of less than a predetermined size are less efficient (see, e. G., "Buergel" p. 141). Therefore, the dislocation density generated through the deformation should not be too high, because this would result in too much but too weak precipitate, and the excess dislocation could move freely, thereby preforming being detrimental. In other words, the maximum deformation is preferably 90% for the hot forming part of 1000 占 폚 or less and 90% for the cold forming.

In the case of molding in the range of 20 to 99%, annealing at 950 캜 to 1050 캜 can restore the tissue. Thereby, the dislocation density is reduced, and thus a positive effect appears again on the distribution of the Laveth phase and / or the Fe-containing particles and / or the Cr-containing particles and / or the Si-containing particles and / or carbides.

The first possibility of introducing the Laveth phase and / or the Fe-containing particles and / or the Cr-containing particles and / or the Si-containing particles and / or carbides into the molded material is to produce the necessary parts as a semi-finished product, Is heated at a rate of 0.1 占 폚 / min to 1000 占 폚 / min to an application temperature of 550 占 폚 to 1000 占 폚. During the heating, the Laves phase and / or the Fe-containing particles and / or the Cr-containing particles and / or the Si-containing particles and / or the carbides are precipitated in such a manner that they are finely distributed in the structure. The fine distribution occurs due to nucleation in the low temperature range and subsequent growth of the nucleus at higher temperatures. Therefore, the heating rate should not be faster than 1000 ° C / min, because otherwise the time for the process becomes too short. Heating rates less than 0.1 ° C / min are uneconomical.

The second possibility is the independent heat treatment of the material. For this purpose, the molded semi-finished product / parts are heat treated in a temperature range of 550 ° C to 1060 ° C under a protective gas or air for a period of from t _min to t _max , and then transferred (blown) Cooled in a gas or air, or in water, or in a furnace for a heat treatment to 800 ° C or below, wherein

및

이고, T_a = T + 273.15이며, t_min 및 t_max의 단위는 분이고 열처리 온도(T)의 단위는 ℃이다. 이 경우 목표하는 부품은 상기 열처리 이전 또는 이후에 제조될 수 있다.

And

T _a = T + 273.15, the units of t _min and t _max are in minutes, and the unit of heat treatment temperature (T) is ° C. In this case, the target part can be manufactured before or after the heat treatment.

It should be noted that for annealing, especially for semi-finished products / parts of thick walls, all parts of the part must reach the minimum annealing time required at a predetermined temperature. This is taken into account when determining the starting point of the annealing time. Similarly, care must be taken that any area of the semi-finished product / part must not exceed the required maximum annealing time.

Times shorter than t _min are not sufficient for the formation of Laveth phase and / or Fe-containing particles and / or Cr-containing particles and / or Si-containing particles and / or carbides. there is a risk that the precipitates become very rough in the case of a time longer than t _max so that the particles can no longer contribute significantly to the creep strength. If the time is longer than t _max , a high temperature of 550 ° C to 1060 ° C is likely to result in recovered tissue, which may still be entirely effective. However, as the recovery progresses, the dislocation density is further reduced, which leads to a gradual non-uniform distribution of precipitates and eventually to a positive effect on creep strength. At low temperatures in the range of 550 ° C to 1060 ° C, the time in excess of t _max is additionally uneconomical.

Annealing can occur under protective gas (a similar atmosphere with reduced argon, hydrogen, and oxygen partial pressures). Cooling is carried out in a protective gas or air which is stationary for economical reasons, in a protective gas or air which moves (blowing) or in air or in water, and furnace cooling should be avoided, especially if the temperature is above 800 ° C. However, if the temperature is less than 800 ° C, furnace cooling is possible.

The chromium content determines the oxidation resistance and thermal expansion coefficient of the material. The oxidation resistance of the material is based on the formation of a closed chromium oxide layer. When the chromium content is less than 12%, an increased iron-containing oxide is produced which, in particular, at relatively high operating temperatures, deteriorates the oxidation resistance. Therefore, the chromium content is adjusted to 12% or more. If the chromium content exceeds 30%, the processability and availability of the material is degraded by the increased formation of the embrittling phase, especially the sigma phase. Therefore, the chromium content is limited to 30% or less. As the chromium content increases, the expansion coefficient decreases.

Accordingly, when applied to a fuel cell, in particular, the expansion coefficient can be set to an area suitable for ceramics in the fuel cell. In this case, the chromium content is about 22 to 23%. However, there are no limitations when applied in other applications, such as reformers or power plants.

When manganese is added, chromium-manganese spinel is formed on the chromium oxide layer formed on the material when the aluminum content is less than 2%. This chromium-manganese spinel reduces chromium evaporation and improves contact resistance. At least 0.001% manganese content is required for this. Manganese in excess of 2.5% forms a very thick chromium-manganese spinel layer and degrades oxidation resistance.

Niobium, molybdenum, tungsten or tantalum may be involved in the formation of precipitates in iron containing alloy, which, for example carbides and / or Laves-phase Fe ₂ (M, Si) or Fe ₇ in the (M, Si) ₆ Lt; / RTI > Molybdenum, tungsten or tantalum is a further solid solution hardener, which contributes to the improvement of creep strength. Here, the lower limit is determined by the fact that some amount must be present for each to be efficient, and the upper limit is determined by the processability. Accordingly, the preferable range is as follows.

Nb: 0.1 - 2%,

W: 0.1 - 5%.

W can be completely or partially replaced by Mo and / or Ta: 0.1 - 5%.

Silicon can be involved in the formation of precipitates in iron-containing alloys, for example in Lavess-like Fe ₂ (M, Si) or Fe ₇ (M, Si) ₆ . Silicon provides increased precipitation and stability on the LABEES, thereby contributing to creep resistance. Upon formation of the Lavess phase, silicon is completely bound in the Lavess phase. Thus, the formation of a silicon oxide layer beneath the chromium oxide layer no longer occurs. At the same time, the incorporation of M in the oxide layer is reduced, thereby preventing a negative influence of M on oxidation resistance. At least 0.05% of Si must be present to achieve the desired action. If the Si content is too high, a negative action of Si may occur again. Therefore, the Si content is limited to 1%.

An aluminum content of less than 1% deteriorates the oxidation resistance because aluminum causes internal oxidation. However, an aluminum content of greater than 1% forms an aluminum oxide layer that does not have electrical conductivity under the chromium oxide layer, thereby reducing contact resistance. Therefore, if the chromium oxide formers are desired or if their oxidation resistance is sufficient, the aluminum content is limited to 0.5% or less. An example of this is an application as an inter-connector plate, for example. However, a predetermined aluminum content of at least 0.001% is required for deoxidation of the melt. If a conductive oxide is not required and at the same time a significantly higher oxidation resistance is required than is provided by the chromium oxide layer, the alloy can form a closed aluminum oxide layer through an aluminum content of at least 2% (DE 101 20 561). An aluminum content of greater than 6.0% results in processing problems and thus the cost increase.

Carbon produces carbide precipitates, which in turn contribute to creep strength. The carbon content should be less than 0.1% to avoid degradation of processability. However, the carbon content should be greater than 0.002% in order to be effective.

The nitrogen content should be at most 0.1% to prevent the formation of nitrides which degrade processability. In addition, the nitrogen content should be greater than 0.002% to ensure the processability of the material.

The sulfur content should be as low as possible because the surfactant element lowers oxidation resistance. Therefore, it is determined as S of 0.01% at maximum.

Oxygen affinity elements such as Ce, La, Pr, Ne, SC, Y, Zr and Hf reduce oxide growth and improve the tackiness of the oxide layer and improve oxidation resistance. It is reasonable that the minimum content of at least one element among Ce, La, Pr, Ne, Sc, Y, Zr and Hf elements is 0.02% in order to obtain the effect of improving oxidation resistance of Y. The upper limit is determined to be 0.3% by weight for cost reasons.

Titanium is bound in the oxide layer at the time of oxidation like each oxygen affinity element. Titanium also causes internal oxidation. However, in this case, the oxides which are generated are small and finely distributed, causing surface hardening, thereby preventing the incorporation of bulges and metal areas of the oxide layer during the oxidation (see DE 10 2006 007 598 A1). The flaking is undesirable because the cracks occurring in this case lead to an increase in the oxidation rate. Accordingly, Ti contributes to improvement of oxidation resistance. With respect to the efficiency of the Ti content, at least 0.01% of Ti should be present, but not more than 0.5%, because if it exceeds 0.5%, it will not improve the effect but increase the cost.

The phosphorus content should be less than 0.030%, because the surfactant element lowers oxidation resistance. Too little P content increases the cost. Therefore, the P content is 0.002% or more.

The magnesium and calcium contents are adjusted to a diffusion range of 0.0001 to 0.05% by weight and 0.0001 to 0.03% by weight, respectively.

It was found that the cobalt content from 3% lowered the oxidation resistance. The lower limit is 0.01 wt.% For reasons of cost. For nickel and copper, the same applies to cobalt.

Boron is limited to a maximum of 0.005%, because the element reduces oxidation resistance.

The contents of the diagram are as follows.
Fig. 1 shows the texture of the hot-formed material in Example 1. Fig.
Fig. 2 is the texture of the hot-formed material after grain annealing at 1075 占 폚 for 20 minutes and cooling in still air in Example 1 (grain size: 137 占 퐉).
Fig. 3 is the texture of the material after annealing at a temperature of 600 DEG C to 1000 DEG C for 20 minutes and cooling in still air in Example 2, respectively.
Fig. 4 is the texture of the material after the annealing at 920 deg. C in air continuously in Example 3, followed by cooling in stationary air for 20 minutes, and cooling in stationary air (etching with V2A caustic).
Fig. 5 is the texture of charge 161061 after cooling in a protective gas flow at low temperature and annealing at 1050 DEG C / 3.4 min under continuous protective gas in Example 4. Fig.
Fig. 6 is a graph showing the results of annealing at 1050 deg. C / 3.4 min under continuous protective gas in Example 4 and cooling in a low temperature protective gas flow and annealing at 1050 deg. C / 20 min in air followed by cooling , And the grain size is 134 mu m (etching with V2A caustic).
Fig. 7 is the texture of the charge 161995 after cooling in a protective gas flow of low temperature and annealing at 1050 DEG C / 3.4 min under continuous protective gas in Example 5. Fig.
FIG. 8 is a graph showing the results of annealing at 1050 DEG C / 3.4 minutes under continuous protective gas in Example 5, cooling in a low temperature protective gas flow, and annealing at 1075 DEG C / And the texture of the charged material 161995, and the grain size is 139 μm.
Fig. 9 is a texture of a hot-formed material after annealing at 1075 占 폚 for 22 minutes and cooling in still air in Example 5, and the grain size is 134 占 퐉 to 162 占 퐉.
10 is a graph showing the results obtained by performing annealing at 1075 캜 for 22 minutes in Example 5 and cooling in air after it was stopped, followed by annealing at 700 캜 for 4 hours, followed by cooling in still air, , And the grain size is 136 mu m.
Fig. 11 is the texture of charge 161995 after cooling in a continuous protective gas at a temperature of 1050 DEG C for 2.8 minutes and a protective gas flow of low temperature.
FIG. 12 shows the results of the annealing at 1050 DEG C for 2.8 minutes under continuous protective gas and cooling in a low temperature protective gas flow, followed by annealing at 1075 DEG C for 10 minutes under air, followed by cooling in still air, And the grain size is 108 mu m.

In the following, objects of the present invention will be described in more detail based on embodiments.

Table 1 shows the analysis of the charge used for the following examples. The charge was melted in an amount of about 30 tons in an arc and then tapped into ladles and then treated with vacuum treatment as well as decarbonization and deoxygenation in a VOD system and cast into blocks. The blocks were then hot rolled and then cold rolled with intermediate annealing according to their respective final thicknesses. After hot rolling, the oxide layer was removed by corrosion.

Materials with the same analytical content as listed in Table 1 primarily precipitate Fe ₂ (M, Si) or Fe ₇ (M, Si) ₆ and a significantly reduced amount of carbides in Laves.

Example 1

In this embodiment, the material consisting of the charge 161061 shown in Table 1 is subjected to a solution heat treatment at a temperature higher than 1070 DEG C over a period of more than 7 minutes and then cooled in a stationary air, , The mechanical shaping was initiated at a starting temperature of greater than 1070 ° C, and the final 78% mechanical strain was performed through rolling at 500 ° C to 1000 ° C.

Fig. 1 shows a typical appearance of the above-described deformed tissue. In the case of the polished surface etched by electrolytic etching using oxalic acid, it can be confirmed that only a slight Laves phase is precipitated to a degree that can be observed with a microscope.

The shaped material is then annealed at 1075 DEG C for 20 minutes and then cooled in still air. This gives a structure containing a slight Lavesite-like precipitate and a grain size of about 137 mu m (Fig. 2), which is a typical coarse grain texture.

The creep test was performed on the material at a temperature of 750 캜 and an initial stress condition of 35 MPa as described above, and the sample was broken after 12.8 hours at an elongation (A) of 69.8% (Table 2). The elongation of the annealed material at room temperature was 35%, which is a very good value for ferrite.

On the contrary, a sample for the creep test is prepared by simulation of the part with the hot rolled material in the same sense as the preform, and then the sample is heated to an application temperature of 750 DEG C at about 60 DEG C / min, The creep test was performed under the initial stress conditions of 35 MPa and temperature. The sample surprisingly fractured after 255 hours at 29% elongation (A), which means that the time to failure was extended by 20 times. The parts can be manufactured very easily because, as described above, the hot-formed state exhibits an elongation of 19% in a room temperature tensile test, which is an excellent numerical value and means good workability of the material.

This example clearly shows that preformed and textured grained structures are better in terms of breakdown time and creep strength, contrary to prior art techniques such as those described in Table 3.7 of "Buergel" 196-199 .

Example 2

In this embodiment, the hot-rolled material according to Example 1 was subjected to annealing in the air at 600 ° C. to 1000 ° C. for 20 minutes, or 240 minutes or 1440 minutes, respectively (t _min according to Equations 1 and 2 and t _max , see Table 3) and then cooled in still air. After the heat treatment, a sample was prepared from the plate and subjected to a creep test at 750 DEG C and 35 MPa stress as described above. The results are summarized in Table 3.

Only a breakdown time of 10.4 h is achieved under the condition that the elongation (A) is 79.5% after 20 minutes at 1000 ° C. At a temperature of 600 ° C to 950 ° C after 20 minutes, a breaking time of more than 100 hours, which is increased by 7 times or more, is achieved under the condition that the elongation (A) is 22.7% or more. The maximum breakdown time for annealing for 20 minutes is achieved at 564 hours at 850 ° C. The maximum breakdown time for annealing for 240 minutes is achieved at 800 ° C and 396 hours. Even after 1440 minutes at 700 ° C, a breaking time of 645h is achieved. Figure 3 shows the structure after various annealing for 20 minutes. The structure of Figure 3 is not globular recrystallized. Tissue shows a typical appearance of deformed tissue up to 850 ° C (maximum value of breakdown time). Recovery from about 900 ° C is evident, but this means that the dislocation density is higher than that of the spherically recrystallized tissue. In the recovered tissue, the dislocations were partially rearranged toward small-angle grain boundaries. This shows an effect similar to preforming. In the case of the polished surface etched by electrolytic etching using oxalic acid, it is observed that the Laves phase is deposited to a degree that can be observed by a microscope from about 750 ° C, and the densely and uniformly precipitated to 850 ° C (maximum value of the fracture time) have. From about 900 ° C, the Laveth phase is clearly identified on the incineration or grain boundaries as well as the precipitates on the grain boundaries, thereby providing incineration or grain boundaries, which corresponds to measure 2 for increasing creep strength See above). At 1000 ℃, it can be seen that very large grains are produced by continuous recovery, so that the dislocation density decreases greatly and the destruction time no longer increases. The maximum value of the breakdown time is provided in a molded tissue comprising a more dense and uniformly precipitated Laves phase.

The plate which has been annealed at room temperature for 20 minutes exhibits an elongation of more than 13% at all temperatures of 600 ° C to 950 ° C, which is regarded as satisfactory for the ferrite alloy and makes the material workable. The elongation is at a minimum in the range of 700 ° C to 800 ° C and is improved toward the annealing temperature at low or normal temperature respectively because the Laves phase has already been precipitated at a low temperature but is not confirmed by the microscope Although they exhibit a lower volume fraction, they are very finely distributed. At higher temperatures, relatively more volume fraction is deposited, but somewhat rougher and with low-angle grain boundaries and grain boundaries.

Annealing at 1000 ° C exceeds t _max = 19.6 minutes with an annealing time of 20 minutes. Thus, the annealing is outside the scope of the present invention and is used only as a reference. Also, the breakdown time is only 10.4 hours. The annealing time for 20 minutes at a temperature of 600 ° C to 950 ° C is within the range according to the invention between t _min and t _max . The breakdown time is therefore clearly increased by at least 7 times as compared to the spherically recrystallized state of coarse grains that occurs after annealing at 1075 DEG C for 20 minutes in Example 1 and subsequent cooling in stationary air, did.

Example 3

In this example, the material consisting of the charge 161061 shown in Table 1 was subjected to a solution heat treatment at above 1070 DEG C over a period of more than 7 minutes, followed by cooling in stationary air and hot rolling to a plate of 12 mm thickness, With a starting temperature exceeding 1070 DEG C, and the last 60% mechanical strain was carried out by rolling at 1000 DEG C to 500 DEG C.

The shaped plate was then annealed continuously at 920 ° C for 28 minutes in the air for industrial use and then cooled in still air. The tensile samples made from the above material thus confirmed to have a fracture time of 391 hours at an elongation (A) of 38% in a creep test at a temperature of 750 DEG C and an initial stress of 35 MPa (Table 4). The tissue was recovered, not recrystallized as spherical. The texture had precipitates in the crystal grains and in each of the incinerator and grain boundaries (Fig. 4). The breakdown time is 30 times the time achieved with a spherically recrystallized coarse grain structure having a grain size of 137 mu m after annealing at 1075 DEG C for 20 minutes. Annealing at 920 占 폚 is within the range according to the invention of t _min = 0.3 min to t _max = 162.6 _min with an annealing time of 28 _min .

The plate thus treated exhibited a very good elongation at room temperature of 18%, an elastic limit of 475 MPa and a tensile strength of 655 MPa (see Table 4), which allows the material to be molded well.

Example 4

In this embodiment, the material consisting of the charge 161061 and the charge 161995 is subjected to a solution heat treatment at a temperature greater than 1070 DEG C for more than 7 minutes, followed by cooling in a protective gas blown, followed by a hot rolling and an oxide layer removal process, Cold-rolled with a 1.5 mm thick plate, where cold forming was carried out at 53%. Followed by annealing at 1050 ° C for 3.4 minutes under a protective gas in a continuous furnace followed by cooling in a low temperature protective gas stream. Charge 161995, as well as charge 161061 (FIG. 5), then showed recovered texture, including elongated grains (FIG. 7) and precipitates on Laves, although less noticeable than in FIG. A portion of the material was then annealed again at 1050 ° C for 20 minutes in air and then cooled in still air. The charge was then recrystallized to a spherical shape, in which case the charge 161061 had a grain size of 134 μm (FIG. 6) and the charge 161,995 had a grain size of 139 μm. Laveth-phase precipitates were only found to be negligible.

Tables 5a and 5b summarize the results of the creep test and the tensile test at room temperature. After annealing at 1050 ° C for 3.4 minutes, charge 161061 exhibited an initial load of 35 MPa and a breakdown time of 25.9 hours at 50% elongation (A) in a creep test at 750 ° C and after additional annealing at 1050 ° C for 20 minutes Produced very coarse grains and the breakdown time was 7.9 hours, decreasing from 1/3 to 83% elongation (A).

Similarly, loading 161,995 exhibited an initial load of 35 MPa and a breakdown time of 33.5 hours at an elongation of 89% at a creep test at 750 ° C and a very rough grain after additional annealing at 1075 ° C for 20 minutes, The breaking time decreased to 1/4 at 92% elongation (A) and was 7.9 hours. In the tensile test at room temperature, elongation was 28% for the charge 161061 and 26% for the charge 161995 under annealing time conditions of 1050 ° C and 3.4 minutes. This is a very good value for ferrite and allows the material to be molded very well. The elongation percentage was 31% for the charge 161061 and 29% for the charge 161995 in the coarse grain structure.

This example shows the effect of annealing time at a temperature of about 1050 < 0 > C. During short-term annealing of a few minutes there is dislocation (deformation) and a sufficient Laves phase in the material, which in this case extends the time from creep test to failure to 3 to 4 times. For longer term annealing, the Laveth phase is fully dissolved as shown in the charge 161061, and the tissue is recrystallized to a spherical shape with a corresponding short time from the creep test to the failure.

Annealing at 1050 ° C for 20 minutes is t _max = 6.0 minutes with an annealing time of 20 minutes. Accordingly, the annealing is not within the scope of the present invention, and is used as a reference for annealing at 1075 캜 for 20 minutes. Annealing at 1050 [deg.] C for 3.4 minutes belongs to the range according to the present invention with tmin = 0.32 min to _tmax = 6.0 _min with an annealing time of 3.4 _min and shows clearly increased breakdown time in the creep test according to the present invention.

Example 5

In this example, the material consisting of the charge 161061 was subjected to a solution heat treatment at above 1070 DEG C for more than 7 minutes, then cooled in still air, and then hot rolled into a 12 mm thick plate. Where the shaping was initiated at a starting temperature of greater than 1070 DEG C and mechanical strain of the last 70% was carried out through rolling at 1000 DEG C to 500 DEG C. [

The shaped material as described above was then heat treated at 1075 DEG C for 22 minutes and cooled in still air, so that a very coarse grained texture with a slight Lavesite precipitate and a grain size of about 134 to 162 mu m was obtained (Fig. 9). The material was subjected to a creep test at an initial stress condition of 40 MPa at a temperature of 700 캜, whereby the sample was destroyed after 228 hours at an elongation (A) of 51% (Table 6). In the creep test at 60 MPa, the sample was destroyed after 8.1 hours at 43% elongation (A). The annealed material at room temperature showed an elongation of 35%, which is a very good value for ferrite.

The material annealed at 1075 ° C for 22 minutes was further annealed at 700 ° C for 4 hours and then cooled in still air. As a result, the Lavess phase was deposited in the tissues (FIG. 10). The creep test was then carried out at an initial stress of 40 MPa at 700 ° C. The sample was destroyed after 104 hours at an elongation (A) of 72.6%. In other words, the sample was destroyed in a much shorter time than after 10 min. When the creep test was carried out at 60 MPa, the sample was broken after 6.3 hours at 63% elongation (A), which was also destroyed in a much shorter time than after 10 min.

This fact is evidence that the precipitation of Lavess phase (s) must be made in tissues with increased dislocation density, that is, in molded or restored tissues, in order to achieve an extension of the disruption time. Precipitation Precipitation into heat-treated tissues causes exactly the opposite, in other words, shortening of the destruction time. The reason for this is in a more uniform distribution of very fine precipitates of calcite in tissues where dislocation densities have risen compared to precipitation in insoluble rough grain grains, i. E. In transformed or recovered tissues.

Example 6

In the present embodiment, the hot-rolled material according to Example 1 was subjected to heat treatment at 750 ° C to 1000 ° C for 20 minutes and at several temperatures for 120 minutes, 240 minutes, 480 minutes, 960 minutes, 1440 minutes, and 5760 minutes (See Table 7 for t _min and t _max according to equations 1 and 2), and then cooled in still air. A sample was prepared from the plate after the heat treatment, and then creep test was performed at 750 占 폚 and under a stress condition of 40 MPa as described above. The reason for selecting the stress higher than in Example 2 was to shorten the test time. The aim here was to identify the time of annealing suitable for annealing. The results are summarized in Table 7.

After 20 minutes at 1000 占 폚 only a breakdown time of 8.8h at an elongation (A) of 78.7% is achieved. In Example 2, a relatively short fracture time was achieved after 20 minutes at 1000 占 폚 and after a creep test at 750 占 폚 and 35 MPa for 20 minutes followed by 1075 占 폚 of solid solution heat treatment followed by quenching in still air , So the above values can only be taken as reference data for the breakdown time in the solid state heat treatment. However, in the case of the variants according to the invention, the breaking time exceeded 1.5 times or more.

After 20 minutes at 750 DEG C to 900 DEG C, the breakdown time is increased by 10 times or more at an elongation (A) of more than 27% to achieve a breakdown time of more than 100 hours. At a 20 minute anneal, the maximum breakdown time is achieved at 850 DEG C and 296 hours. At 120 min of annealing, the maximum breakdown time is achieved at 800 DEG C and 227 hours. At 240 min annealing, the maximum breakdown time is achieved at 750 ° C for 182 hours, and there is no numerical value for 700 ° C. At 480 min annealing, the maximum breakdown time is achieved at 800 ° C and 169 hours. In the case of 960 minutes, only the breakdown time for 750 ℃ could be determined as 139 hours at 24.2% elongation. After 1440 minutes and 5760 minutes, only a fracture time which is clearly less than the maximum fracture time achieved at this temperature at 750 ° C and 800 ° C is achieved. For example, after a treatment time of 480 minutes at 800 ° C, the breakdown time of 169 hours is reduced to a value of 46 hours after the treatment time of 1440 minutes and to a value of 17.5 hours after the treatment time of 5760 minutes, Are within the scope of the present invention. A further increase in treatment time further reduces the breakdown time, meaning that the t _max of 7059 minutes somewhat exceeds the time of 5705 minutes. All stretching rates for the heat treatment temperature of 750 ° C to 900 ° C and for a time of 20 minutes to 5760 minutes are 24.2% to 43%, and thus greater than 18% as required to prevent brittle fracture. In the case of tissues after 20 minutes of annealing, the matters mentioned in Example 2 are applied because the annealing method is the same. In addition, even at a relatively higher stress of 40 MPa in the creep test, the maximum value of the fracture time is provided in a modified structure comprising a more dense and uniformly precipitated Laves phase.

The plate annealed at room temperature for 20 minutes exhibits an elongation of more than 13% at all temperatures between 600 ° C and 900 ° C as in Example 2, which is also regarded as satisfactory for the ferrite alloy and makes the material workable.

Example 7

In this example, a 1.5 mm thick loading 161995 material was used, which was cold rolled at a temperature of 1050 ° C for 3.4 minutes under a protective gas in a continuous furnace at 53%, cooled in a low temperature protective gas stream and then annealed. In the same manner and manner, 40% cold forming at 10.degree. C. for 2.8 minutes under a protective gas in a continuous furnace followed by cooling in a low temperature protective gas stream followed by annealing to obtain the material of 161,995 as a 2.5 mm thick . Thereafter, the 2.5 mm thick material, as evident from FIG. 4, although less obviously, is formed from the restored texture (FIG. 11) containing elongated grains as in FIG. 7 and the material according to Example 4 and the Lavess- Respectively. A portion of the material was then annealed again at 1050 < 0 > C for 10 minutes under air and then cooled in still air. The material was then recrystallized to a spherical shape with a grain size of 108 mu m. Laveth-phase precipitates were only found to be negligible. Thereafter, the material treated at 1050 DEG C / 2.8 minutes and the material treated at 1075 DEG C / 10 minutes as the final heat treatment were rolled to a molding of 2.8 to 40%. Then creep tests at 750 ° C and 35 MPa and tensile tests at room temperature were carried out. The results are summarized in Table 8.

In an creep test with an initial stress of 35 MPa at 750 ° C after 3.4 minutes of annealing at 1050 ° C, loading 161995 exhibited a breakdown time of 33.5 hours at an elongation of 89% (A) and after additional annealing at 1050 ° C for 10 minutes It produced very coarse grains and its breakdown time decreased to 1/3 from 50.4% elongation (A) to 10.8 hours.

A sample for the creep test was prepared as a simulation for the part with the material molded according to the conditions of 1050 DEG C / 2.8 min, and then the sample was heated to an application temperature of 750 DEG C at about 60 DEG C / min. The creep test was carried out at an initial stress of 35 MPa at a temperature of 750 ° C. As a result, the fracture elongation at the time of molding of 5 to 40% decreased to a value of about 10 hours at a fracture elongation of more than 45%.

On the contrary, a sample for the creep test was prepared as a simulation for the part with molded material at 1050 ° C / 10 minutes, and then the sample was heated to an application temperature of 750 ° C at about 60 ° C / min. The creep test was carried out at an initial stress of 35 MPa at a temperature of 750 ° C. Accordingly, the fracture elongation at 2.9 to 40% during molding increased to a value from 49 to 137 hours, which means that the fracture time increased by more than 4 times as compared to the material molded at 1050 DEG C / 2.8 min, And the fracture elongation is 18.9 to 60%.

However, from 10% molding, the tensile test at room temperature showed a fracture elongation of less than 8%. In other words, the preferable degree of molding for cold forming is 0.05 to 10%. Accordingly, in this embodiment, it was confirmed that the annealing before molding (in this case, at 1050 ° C for 2.8 minutes) at a too low temperature and a too short time did not increase the fracture elongation after molding. The increase in breakdown elongation was achieved after annealing over 1050 ° C for a period of more than 6 minutes (in this case 7 minutes at 1075 ° C).

The description / explanation of the table is as follows.

Table 1 shows the composition of the analyzed alloys (all indications are% by weight).

Table 2 shows the results of a creep test at 750 DEG C and 35 MPa and a tensile test at room temperature in relation to the hot rolling and heat treatment according to Example 1 for a 12 mm thick plate (R: Reference material, E: results according to the present invention).

Table 3 shows the results of the creep test at 750 DEG C and 35 MPa and the tensile test at room temperature in relation to the hot rolling according to Example 1 for the 12 mm thick plate and the heat treatment according to Example 2 (R: , E: result according to the present invention).

Table 4 shows the results of a creep test at 750 DEG C and 35 MPa and a tensile test at room temperature conducted in relation to Example 3 on a 12 mm thick plate (R: reference material according to prior art, E: result).

Table 5 shows the results of a creep test at 750 DEG C and 35 MPa and a tensile test at room temperature conducted in relation to Example 4 on a strip of 1.5 mm thickness (R: reference document according to the prior art, E: result).

Table 6 shows the results of the creep test at 750 DEG C and the tensile test at room temperature conducted in relation to Example 5 on a 12 mm thick plate (R: reference material according to prior art, E: results according to the present invention) .

Table 7 shows the results of a creep test at 750 DEG C and 40 MPa and a tensile test at room temperature in relation to hot rolling and heat treatment in Example 6 for a 12 mm thick plate (R: reference material according to prior art, E : Results according to the present invention).

Table 8 shows the results of a creep test at 750 DEG C and 35 MPa and a tensile test at room temperature conducted in relation to Example 7 for strips 1.5 mm to 2.5 mm thick consisting of the charge 161995 (R: Data, E: results according to the present invention).

[Table 1]

[Table 2]

[Table 3]

[Table 4]

[Table 5a]

[Table 5b]

[Table 6]

[Table 7a]

[Table 7b]

[Table 8a]

[Table 8b]

Laves Phase: Laves Phase

Claims (50)

A semi-finished product made of an Fe-containing particle, a Cr-containing particle, a Si-containing particle and an iron-chromium alloy for precipitating at least one of carbides is subjected to a thermomechanical treatment to form a part made of the alloy , Comprising:
In the first step, the alloy is subjected to a solution heat treatment at a temperature above the solidification heat treatment temperature and then cooled in a protective gas or air that is moving or blown in a stationary protective gas or air,
In the second stage, the semi-finished product is subjected to mechanical forming in the range of 0.05 to 99%, and in the subsequent stage, the semi-finished product is heated at a heating rate of 0.1 ° C / min to 1000 ° C / min by applying load to the application temperature of 550 ℃ to 1000 ℃, Laves phase Fe ₂ (M, Si), Laves phase Fe ₇ (M, Si) _6, Fe-containing particles, Cr-containing particles, Si-containing particles, and the carbide At least one of which is precipitated in a dispersed form,
The alloy has the following composition (in weight%) and is given mechanical deformation at room temperature, measured in plastic elongation in tensile tests of greater than 13%:
Cr 12 - 30%
Mn 0.001 - 2.5%
0.1 to 2% of Nb,
W 0.1 - 5%,
0.05 - 1% of Si,
C 0.002 - 0.1%
N 0.002 - 0.1%
S Up to 0.01%
Fe rest, and
Unavoidable molten impurities. A semi-finished product made of an Fe-containing particle, a Cr-containing particle, a Si-containing particle and an iron-chromium alloy for precipitating at least one of carbides is subjected to a thermomechanical treatment to form a part made of the alloy , Comprising:
In the first step, the alloy is subjected to a solution heat treatment at a temperature above the solidification heat treatment temperature and then cooled in a protective gas or air that is moving or blown in a stationary protective gas or air,
The mechanical molding of the semi-finished product in the range of 0.05 to 99% in step 2 (mechanical forming) is performed and, and under such a molded semi-finished product protection for a period of time t _min to t _max gas or air in a subsequent step 550 to 1060 ℃ And then cooling in water or in protective gas or air moving or blown in stationary protective gas or air or cooling to below 800 ° C in a furnace for heat treatment, Beth phase Fe ₂ (M, Si), Laves phase Fe ₇ (M, Si) _6, Fe-containing particles, Cr-containing particles, Si-containing particles and are deposited in the dispersed form at least one of a carbide, wherein the parts are Finished before or after the heat treatment,
Where t _min and t _max are calculated according to the following formula,

And

, T _a = T + 273.15,
The time t _min and t _max are in minutes, the unit of heat treatment temperature (T) is in ° C,
The alloy has the following composition (in weight%) and is given mechanical deformation at room temperature, measured in plastic elongation in tensile tests of greater than 13%:
Cr 12 - 30%
Mn 0.001 - 2.5%
0.1 to 2% of Nb,
W 0.1 - 5%,
0.05 - 1% of Si,
C 0.002 - 0.1%
N 0.002 - 0.1%
S Up to 0.01%
Fe rest, and
Unavoidable molten impurities. 3. A method according to claim 1 or 2, wherein in said first step said alloy is subjected to a solution heat treatment at a temperature of at least 1050 DEG C for a period of time greater than 6 minutes and less than 24 hours. 3. A method according to claim 1 or 2, wherein in said first step said alloy is subjected to a solution heat treatment at a temperature of at least 1060 DEG C for a period of time greater than 1 minute and less than 24 hours. delete The process according to any one of claims 1 to 3, wherein after cooling in a protective gas or air, which is moved or blown in a stationary protective gas or air, or in water, at a temperature above the solidification heat treatment temperature, (M, Si) ₆ , Fe-containing particles, Cr-containing particles, Si-containing particles, and carbides in the Laves phase Fe ₂ (M, Si), Laves phase Fe ₇ Wherein at least one of the first and second layers is present or does not exist at all. 7. The method according to claim 6, wherein the solid solution heat treatment is performed at a temperature of 1050 DEG C or more for each of more than 6 minutes and less than 24 hours, or 1060 DEG C or more for more than 1 minute and less than 24 hours. The component manufacturing method according to claim 1 or 2, wherein the molding of the semi-finished product is performed by hot forming. Process according to claim 1 or 2, characterized in that the hot forming of the semi-finished product starts at a starting temperature of more than 1070 DEG C and the last mechanical working of 0.05 to 90% is carried out at 1000 to 500 DEG C. . Process according to claim 1 or 2, characterized in that the hot forming of the semi-finished product starts at a starting temperature of more than 1070 DEG C and the last mechanical working of 0.05 to 95% is carried out at 1000 to 500 DEG C. . Process according to claim 1 or 2, characterized in that the hot forming of the semi-finished product starts at a starting temperature of more than 1070 DEG C and the last mechanical working of 0.05 to 90% is carried out at 1000 to 500 DEG C. . The component manufacturing method according to claim 1 or 2, wherein the hot forming of the semi-finished product is followed by cold forming. The component manufacturing method according to claim 1 or 2, wherein the molding of the semi-finished product is performed by cold forming. 14. The component manufacturing method according to claim 13, wherein the degree of cold forming of the semi-finished product is 0.05 to 99%. The component manufacturing method according to claim 13, wherein the cold forming of the semi-finished product is 0.05 to 95%. 14. The component manufacturing method according to claim 13, wherein the cold forming of the semi-finished product is 0.05 to 90%. 3. A method according to claim 1 or 2, wherein the mechanical forming of the semi-finished product is 20 to 99%, after which the molded semi-finished product has a temperature range of 950 ℃ to 1060 ℃ under protection for a period of t _min to t _max gas or air And then cooled in protective gas or air, which is heat-treated in a protective gas or air, which is then moved or blown, or in water, after which the part is completed.
here

And

T _a = T + 273.15, the units of t _min and t _max are in minutes, and the unit of heat treatment temperature (T) is ° C. 3. The method according to claim 1 or 2, wherein the alloy further comprises 0.02 to 0.3% (by weight) of La. 3. The method according to claim 1 or 2, wherein the alloy further comprises 0.01 to 0.5% (by weight) of Ti. The method according to claim 1 or 2, wherein the alloy further comprises 0.02 to 0.3% of at least one of Ce, Pr, Ne, Sc, Y, Zr or Hf elements. 3. A method according to claim 1 or 2, wherein the alloy further comprises 0.001 to 0.5% (by weight) of Al. 3. A method according to claim 1 or 2, wherein the alloy further comprises 2.0 to 6.0% (by weight) of Al. 23. The method of claim 22, wherein the alloy further comprises (by weight percent) 2.5 to 5.0% Al. 3. The method according to claim 1 or 2, wherein the alloy further comprises 0.0001 to 0.07% of Mg, 0.0001 to 0.07% of Ca, and 0.002 to 0.03% of at least one P element. The component manufacturing method according to claim 1 or 2, wherein the alloy further comprises 0.01 to 3.0% of at least one of Ni, Co or Cu elements. 3. A method according to claim 1 or 2, wherein the alloy further comprises 0.005% or less of B. The iron-chromium alloy according to claim 1 or 2, wherein the iron-chromium alloy treated with the processing heat treatment and having a Lavess phase dispersed therein contains the following composition (by weight%):
Cr 12 - 30%
Mn 0.001 - 2.5%
0.1 to 2% of Nb,
W 0.1 - 5%,
0.05 - 1% of Si,
C 0.002 - 0.03%
N 0.002 - 0.03%
S Up to 0.005%
Fe rest, and
Unavoidable molten impurities. 3. The method according to claim 1 or 2, wherein the alloy further comprises 0.02 to 0.2% (by weight) of La element. 3. A method according to claim 1 or 2, wherein the alloy further comprises 0.02 to 0.2% (by weight) of Ti. The method according to claim 1 or 2, wherein the alloy further comprises 0.02 - 0.2% (by weight) of at least one of Ce, Pr, Ne, Sc, Y, Zr or Hf elements. The method of claim 1 or 2, wherein the alloy further comprises 0.0001-0.05% by weight of Mg, 0.0001-0.03% Ca, 0.002-0.03% of at least one P element Method of manufacturing parts. 3. A method according to claim 1 or 2, wherein the alloy further comprises 0.003% (by weight) of B or less. The component manufacturing method according to claim 1 or 2, wherein the Nb content (by weight%) is 0.3 to 1.0% and the Si content is 0.15 to 0.5%. 3. A method according to claim 1 or 2, wherein the W content is completely or partially replaced by at least one of Mo and Ta elements. 3. The method according to claim 1 or 2, wherein the alloy comprises (by weight%) at most 0.2% V and at most 0.005% S. 3. A method according to claim 1 or 2, wherein the alloy comprises up to 0.01% O by weight (by weight). 3. The method according to claim 1 or 2, wherein the alloy contains at most 0.01% (by weight) of Zn, Sn, Pb, Se, Te, Bi and Sb elements. 3. A method according to claim 1 or 2, wherein the semi-finished product is formed of a plate, a strip, a rod, a forgings, a pipe or a wire. The component manufacturing method according to claim 1 or 2, wherein the heat treatment is performed for the first time after completion of the component. A high creep strength is produced in at least one of said semi-finished products and said parts having an elongation percentage exceeding 13% by room temperature tensile test by the processing heat treatment of said semi-finished product Gt; As an article,
(By weight) of the following chemical composition,
Cr 12 - 30%
Mn 0.001 - 2.5%
0.1 to 2% of Nb,
W 0.1 - 5%,
0.05 - 1% of Si,
C 0.002 - 0.1%
N 0.002 - 0.1%
S Up to 0.01%
Fe rest, and
Unavoidable molten impurities,
After the processing heat treatment, the Laveth phase (s) exhibits a deformed tissue structure that is einlagerung dispersed in the tissue dislocation of the tissue structure, and creep using 35 MPa at an elongation of 750 ° C and 18% Wherein the breaking time in the test is greater than 1.5 times the breaking time of the grained fully recrystallized tissue within the tissue structure. As an article,
(By weight) of the following chemical composition,
Cr 12 - 30%
Mn 0.001 - 2.5%
0.1 to 2% of Nb,
W 0.1 - 5%,
0.05 - 1% of Si,
C 0.002 - 0.1%
N 0.002 - 0.1%
S Up to 0.01%
Fe rest, and
Unavoidable molten impurities,
After the processing heat treatment, the Lavess phase (s) exhibit a modified tissue structure that is einlagerung in a dispersed form within the tissue dislocation of the tissue structure,
The breaking time (h) exceeding 3 times the breaking time of grained fully recrystallized tissue in a creep test at 35 MPa at 750 DEG C and an elongation of 18% or more is set within the tissue structure. An interconnection in a fuel cell comprising a component manufactured according to claim 1 or 2. A material for a reformer of a fuel cell comprising a part manufactured according to claim 1 or 2. An exhaust gas conduit of an internal combustion engine comprising a component manufactured according to claim 1 or 2. A steam boiler comprising parts made according to paragraph 1 or 2. 43. The article of claim 41 or 42, wherein the article is a metal part or semi-finished article. A material of a heat exchanger of a fuel cell comprising a part manufactured according to claim 1 or 2. A superheater comprising parts manufactured in accordance with claims 1 or 2. 6. A turbine comprising parts manufactured according to claim 1 or 2.

Images