JP6515276B2 - High strength ferritic heat resistant steel structure and method of manufacturing the same - Google Patents

Description

  The present invention relates to a high-strength ferritic heat-resistant steel structure and a method for producing the same, for example, a steel structure subjected to stress at high temperatures for a long period of time, in particular, high for use in members such as power plants and chemical plants. The present invention relates to a welded joint of high strength ferritic heat resistant steel and a method of manufacturing the same.

The prevention of global warming is an urgent issue, and technological development to effectively utilize energy resources is extremely important. In particular, among energy resources, power plants that convert fossil fuels and nuclear fuels into electrical energy, specifically, coal-fired power plants, natural gas direct-fired thermal power plants, and nuclear power plants, have problems with the life of energy resources. There is a need to further improve power generation efficiency (energy conversion efficiency). In addition, improvement of production efficiency is also required for petroleum refining plants and coal gasification plants.
However, at present, for example, the efficiency of thermal power generation has stagnated at 40 to 50%, and in order to suppress the increase of carbon dioxide emissions, further higher efficiency of power generation is required.

  Moreover, the thermal efficiency in each plant is substantially determined by the operating temperature and pressure of the plant as well as the above-mentioned power plant, and particularly in the power plant, the higher the temperature of the steam that drives the turbine of the generator, Conversion efficiency increases.

  At present, the steam temperature of the coal-fired power plant is 620 ° C. However, if the steam temperature is increased by 100 ° C., the efficiency can be expected to be improved by about 5% and by 200 ° C. by about 10%. However, in order to raise the temperature of the steam that drives the generator's turbine, it is necessary to improve the performance of heat resistant steel used not only in the turbine components but also in the heat exchangers and piping.

Of the performance requirements for heat resistant steels, creep properties are important and require long-term creep rupture so that the plant can be operated for several decades.
As a heat resistant steel with improved cream properties, research and development of 9% Cr ferritic heat resistant steel has been conducted up to this point with 600 ° C. as the upper limit of operating temperature. High-temperature ferritic heat-resistant steels such as STBA 29 have been developed and put to practical use. These ferritic heat-resistant steels have a low coefficient of thermal expansion, have resistance to creep-fatigue failure due to thermal stress and deformation of piping members, and are characterized in that they have the same weldability and workability as ordinary steel materials. In addition, the alloy content of these ferritic heat resistant steels is smaller than that of austenitic heat resistant steels used at higher temperatures, and the point of being excellent in economy as well is attractive from an industrial viewpoint.

However, since the atomic structure of iron in these ferritic heat resistant steels is BCC (body-centered lattice), the lattice constant is large, and the strength is low at high temperature. Further, since the atomic diffusion in the steel is faster as compared with the austenitic heat-resistant steel, the structure recovery is also faster, so that the creep strength is lower than the austenitic heat-resistant steel, which is a problem as material characteristics.
From these reasons, in recent years, expectations for the development of a ferritic heat-resistant steel having high creep rupture strength are increasing, and development of a ferritic heat-resistant steel replacing the austenitic heat-resistant steel is being promoted.

  The metallographic structure of ferritic heat resistant steel is martensite or bainite with high dislocation density, and in order to improve the creep characteristics at high temperatures exceeding 600 ° C, it is important to suppress creep recovery of the structure with high dislocation density. Become. In such a heat resistant ferritic steel, although the diffusion of atoms in the steel is fast at high temperatures, the precipitates in the steel effectively act to suppress the movement of dislocations. Therefore, for the purpose of effectively suppressing dislocation movement at high temperatures, that is, creep deformation, in ferritic heat resistant steel, establishment of component design and manufacturing method for introducing stable precipitates into steel is important. It has been pursued.

However, when such heat resistance steel is considered as a component of a power plant, most of these technological developments have been carried out mainly for the purpose of improving creep characteristics as a so-called "base material", and creep strength improvement of "welded joints" There are not many things to be aware of.
For example, although the allowable stress of a 9% Cr high strength ferritic heat resistant steel such as the above-mentioned fire STBA / STPA 28 or fire STBA29 / Fire STPA 29 is specified in the standard, the allowable stress is approximately the creep characteristics and high temperature strength of the base material. It is decided.

  A heat resistant ferritic steel has a so-called transformation point at which temperature transformation occurs between a stable α phase at room temperature and a stable γ phase at high temperature when a low temperature transformation structure (mainly bainite or martensite) is used as the starting structure. . This transformation point contributes to the formation of a high strength, low temperature transformation structure including high density dislocations. However, since the transformation point itself is accompanied by a large change in the structure of the steel (reordering of atoms forming the crystal lattice), the structure of the heat-resistant steel subjected to the thermal history across the transformation point originally has high creep strength. It will be very different from the initial tempering organization introduced to give.

The structure of the heat-affected zone (hereinafter referred to as HAZ in the present invention) of the welded joint is most strongly affected by this phenomenon. The fusion of HAZ and weld metal is at a high temperature of 1500 ° C or higher, and when there is a thermal effect from here toward the base material, it changes with the maximum achieved temperature (maximum heating temperature) for each part according to the distance from the weld metal Form a continuum of That is, HAZ has a texture such that the metallographic structure produced when the maximum heating temperature changes from room temperature to 1500 ° C. is continuous depending on the distance from the weld metal. However, this structure is characterized by the fact that the holding time at the maximum heating temperature is as short as a few seconds, and from the side close to the weld metal, "coarse-grained HAZ (hereinafter, also referred to as coarse-grained area or fine-grained HAZ area)" , "Fine-grained HAZ (hereinafter, also referred to as fine-grained area)", "two-phase area HAZ (hereinafter, also referred to as two-phase area)".
FIG. 1 shows each part including the HAZ of the welded joint and the tissue configuration according to this classification. As shown in FIG. 1, HAZ is formed between the weld metal 1 and the base metal 5, and this HAZ is in the order of coarse grain area HAZ2, fine grain area HAZ3, and two phase area HAZ4 in this order from the weld metal 1 side. It consists of

Among various parts of HAZ, creep damage occurs in the “fine grain area”, and the phenomenon of fracture from the welded joint is called “Type IV damage”, and the Type IV damage occurs in the ferritic heat resistant steel structure. It has not been solved yet, and the solution is regarded as a problem in recent years. That is, Type IV damage is a fracture phenomenon unique to welded joints, in which only the welded joints creep and result in failure despite the base material being in a creep environment, and the temperature and time conditions being able to be used properly. is there.
Recently, it has been empirically found that it is inevitable that this phenomenon occurs in conventional materials (materials that have already been registered in standards and whose allowable stress has been determined). From that, it is in a situation where a safety factor of creep strength, which is called “weld joint creep strength reduction factor” for safe operation including welding joints, is proposed.

Unexamined-Japanese-Patent No. 2008-214753 JP 2008-248385 A JP 2008-266785 A JP 2008-266786 A JP, 2009-293063, A JP, 2010-7094, A JP, 2001-003120, A

This Type IV damage is generated and unavoidable in all of the heat resistant ferritic steels that have been put into practical use, and the generation mechanism has been variously discussed.
The area that becomes fine grain HAZ before welding is a ferritic heat-resistant steel that originally has the same structure as the base material, but this area is heat of HAZ that is exposed to the temperature immediately above the transformation point for several seconds by welding. Receive a cycle. Also, due to the heat from welding, the matrix itself undergoes α → γ transformation, and due to the difference in solid solution limit of α phase and γ phase C, carbides originally precipitated coarsely (mainly M ₂₃ for ferritic heat resistant steels) Carbides other than C ₆ -type carbides are instantly redissolved in the γ phase. However, particularly in fine-grained HAZ, it is thought that the largest cause of Type IV damage is that several tens of percent of coarsely precipitated carbides are shrunk but remain undissolved.
Usually, the welded joint is subjected to post-weld heat treatment (also called stress relief annealing, SR treatment). If the heat treatment temperature is a high temperature different from the tempering temperature by only a few tens of degrees, the carbide remaining in the solid solution state as described above becomes a new precipitation nucleus of the carbide forming element together with the solid solution carbon, resulting in heat At the same time as the cycle coarsens, it reduces the chance of precipitation of fine carbides. That is, the inventors of the present invention have studied that coarse carbides which were precipitated before welding remain undissolved, because the so-called carbides cause loss of "precipitate strengthening ability". The results were found.

  Therefore, it can be understood that Type IV damage is unavoidable in a heat-resistant steel which has the above-mentioned transformation point and in which the creep rupture strength is enhanced by containing carbon and precipitating carbide. That is, Type IV damage is a heat-resistant steel that can be creep-reinforced using carbide, and can occur with any steel type, and is a phenomenon that occurs with low Cr heat-resistant steel without being limited to high Cr and high. In other words, it can be said that it is extremely difficult to prevent Type IV damage with a ferritic heat resistant steel, since it is safe to say that there is no ferritic heat resistant steel that does not apply precipitation strengthening ability by carbide to creep strengthening.

As a cause of Type IV damage, conventionally, Type IV damage is considered to be "softening of strength in fine grained area" because it occurs in fine grained area, or "deterioration of hardenability" caused by small crystal grain size There was a time when it was thought that it was also a loss of dislocation strengthening. However, it has become clear as a result of detailed research that the “fine grained area” has higher room temperature strength than the “two-phase area” or “two-phase area” near base material, and these hypotheses are currently supported Not. Moreover, although no confirmation has been obtained as to the presence or absence of “deterioration in hardenability”, it is generally understood that the strengthening factor that governs the long-term creep strength is not dislocation but precipitation. The ground mechanism is not shown for the creep strength reduction mechanism due to the fine grain structure. That is, these hypotheses have not been able to determine the direct cause of Type IV injury.
Furthermore, it is known experimentally that the relationship between creep strength and grain size is inversely proportional to austenitic heat resistant steels that can only deform grain boundaries, but correlation is made to ferrite based heat resistant steels that can be deformed uniformly. It is known that there is no. Therefore, even if it is possible to create a weld heat-affected zone that does not generate a "fine-grained area" or to display a ferritic heat-resistant steel where a "fine-grained area" is unlikely to occur, coarsening of carbides by the HAZ thermal cycle must be prevented. It came to the conclusion that it is difficult to completely prevent Type IV injury.

In Patent Documents 1 to 4, heat treatment prior to welding is applied to the entire steel pipe containing 50 ppm or less of B for the purpose of preventing the formation of such fine grained area with the conventional ferritic heat resistant steel. Technology to perform type IV treatment and make it possible to prevent Type IV damage.
The technique leaves residual .gamma. At martensite laths or bainite grain boundaries by a short time normalizing treatment before welding to promote their growth and coalescence in reheating at welding, and high temperatures in the base metal before welding It is a technology that makes use of the "tissue memory effect" that reproduces the old gamma particles that were generated in.
In this technology, the prevention of the formation of fine grain areas is completely achieved, and since the carbide is completely dissolved by short-time normalizing and treatment, it is a technology that can completely prevent Type IV damage. However, since the furnace for heat-processing high temperature to the whole member (a steel pipe more than 10 m in length in most cases) containing a groove before welding is needed, construction in the field is difficult. Furthermore, the risk of deformation of the product steel pipe by heating the entire steel pipe, the time for reheating, and the large process load become problems, and it has not become a practical solution for on-site construction .

On the other hand, Patent Documents 5 and 6 disclose the proposal of a steel pipe using a steel material component which does not require the growth and coalescence of residual γ while having the same “texture memory effect (hereinafter, also simply referred to as memory effect)”. .
However, the techniques described in Patent Documents 5 and 6 are known as techniques for exerting a memory effect of shear type α → γ transformation which is generated by adding B at a high concentration exceeding 100 ppm. Further, the techniques described in Patent Documents 5 and 6 are the same as the techniques described in Patent Documents 1 to 4 described above in that old gamma grains of the base material are reproduced at high temperatures, For some reason, it has been considered not to cause Type IV damage.

  However, in the case of a high B content steel as described in Patent Documents 5 and 6, although "the prevention of formation of fine grained area" is achieved, in the area equivalent to the fine grained area, it is accompanied by the dissolution of carbide in a short time. "Coarsening of carbides" through partial solution and reprecipitation can not be avoided. The techniques described in Patent Documents 5 and 6 completely recrystallize to form a fine grain region in that the crystal structure is equivalent to that of the base material and the precipitation position of the carbide remains the large angle grain boundary by the memory effect. In addition, the effect of reducing (delaying) the occurrence of Type IV damage is observed when compared to the conventional ferritic heat-resistant steels in which the precipitation position of the coarsening carbide is unrelated to the newly formed grain boundaries. . Therefore, in the case of a high B-containing steel as described in Patent Documents 5 and 6, Type IV damage does not occur in tests up to several tens of thousands of hours, and creep strength equivalent to that of the base material can be obtained. However, in a long-term creep environment, the decrease in creep strength of HAZ can not be avoided because the coarsening of the carbide precedes, and as a result of the long-term creep test, particularly, it is more than 30,000 hours It became clear from the creep test results.

  In addition, the technique of making the structure equivalent to a base material by the reheat processing (normalization-tempering) of the welded steel pipe is described in patent document 7 other than these methods. Although this method is also aimed at eliminating the uneven strength of the joint by heat treatment including the weld metal, a furnace larger than the heat treatment furnace described in Patent Document 1 is required, and The workability is low. Furthermore, it is practically impossible to reheat all the structures after welding together, especially renormalization, and it is inevitable that a huge cost increase will be inevitable, so it will be a partial heat treatment by all means, and must always be Type IV damage. In the boiler, for example, there are several places where a portion of That is, the technology as described in Patent Document 7 requiring a large-sized furnace can not but be said to be incomplete as a technology for countermeasure against Type IV damage.

  That is, a ferritic heat-resistant steel structure having a low temperature transformation structure which is excellent in on-site workability and does not completely cause Type IV damage has not been developed, and no technology for preventing Type IV damage has been proposed so far. The

  The present invention has been made in view of such circumstances, and has high creep properties and toughness that enables reliable prevention of Type IV damage that occurs in the weld joint heat affected zone (HAZ). It is an object of the present invention to provide a high strength ferritic heat resistant steel structure and a method of manufacturing the same.

  The present invention has been made based on the above-mentioned new findings by the present inventors, and the gist thereof is as follows.

[1] A high-strength ferritic heat-resistant steel structure comprising a base material having a thickness of 5 mm or more, a welding heat affected zone, and a weld metal,
The composition of the base material is, by mass%, C: 0.01 to 0.12,
Si: 0.02 to 0.45,
Mn: 0.20 to 0.60,
Cr: 8.0 to 12.0,
W: 2.00 to 7.50,
Nb: 0.02 to 0.10,
Ta: 0.10 to 1.00,
V: 0.10 to 0.50,
N: 0.003 to 0.015,
B: 0.008 to 0.050,
Co: 0.50 to 3.00,
Nd: 0.005 to 0.05
Contains
Mo ≦ 0.030,
Ni ≦ 0.10,
Cu <0.05,
Al ≦ 0.005,
P ≦ 0.020,
S ≦ 0.010,
O ≦ 0.010
And the balance consists of Fe and unavoidable impurities ,
Before SL (the M and total Cr, Fe, and W, 70% or more) M ₂₃ C ₆ type carbides precipitated in large angle grain boundaries on the weld heat affected zone and the average value of equivalent circle diameter of the TaC is 300nm or less the large angle surface distance of M ₂₃ C ₆ type carbides and TaC precipitated in the grain boundary on the at 150nm or less in the Ku H AZ relation to the particle type, the large angle grain by the M ₂₃ C ₆ type carbides and the TaC A high-strength ferritic heat-resistant steel structure characterized in that the area coverage is at least 40%.
[2] Furthermore, the composition of the base material is, by mass%,
Ti: 0.005 to 0.15%,
Zr: 0.005 to 0.15%
The high-strength ferritic heat-resistant steel structure according to the above [1], which contains one or two of them.
[3] Furthermore, the composition of the base material is, by mass%,
Ca: 0.0003 to 0.0050%,
Mg: 0.0003 to 0.0050%,
Y: 0.0100 to 0.0050%,
Ce: 0.0100 to 0.0050%,
La: 0.0100 to 0.0050%
Among the above, the high-strength ferritic heat-resistant steel structure according to the above [1] or [2], which contains one or more of them.

[4] A method for producing the high strength ferritic heat resistant steel structure according to any one of the above [1] to [3],
The steel piece of the composition according to any one of the above [1] to [3] is heated to 900 ° C. to 1100 ° C. and hot rolled to form a steel plate having a thickness of 5 mm or more,
Then, the steel sheet is subjected to groove processing, and a portion of 20 mm or more and 100 mm or less in the direction orthogonal to the weld line from the position where the groove surface contacts the base material surface, and the width or length of the heat resistant steel structure The portion which is 50% or less of the above is locally heated and maintained at a temperature of 1000 ° C. or more and 1100 ° C. or less so as to be a heat treatment time Ts or more for 30 mm or more determined by the following equation (1) And let it cool down,
Next, after welding the groove, the region between 20 mm and 100 mm is heated to a temperature zone of 720 ° C. or more and 760 ° C. or less in the direction perpendicular to the weld line from the position where the groove face contacts the base material surface. A method for producing a high-strength ferritic heat-resistant steel structure, characterized in that the temperature is maintained for 1 hour or more and then allowed to cool.
Ts = 30 × (t / 25) ^1.2 +26 (1)
However, t is a plate thickness (mm) of a base material, and Ts is a total (minute) of heating time at a heating set temperature (1000 ° C. to 1100 ° C.) of a groove before welding. In addition, t shall be 5 mm or more.

  According to the present invention, it is possible to provide a high-strength ferritic heat-resistant steel structure having excellent creep properties and toughness, and a method of manufacturing the same, while reliably preventing Type IV damage that reduces creep rupture strength of welded joints. In particular, according to the present invention, a ferritic heat-resistant steel structure having a creep rupture strength equivalent to that of a base material for 100,000 hours can be obtained.

FIG. 1 is a schematic cross-sectional schematic view of each portion including a heat-affected zone of a welded joint and a joint for describing each structure. FIG. 2 is a schematic cross-sectional view in the vicinity of a large angle grain boundary for explaining a structure model of a weld heat affected zone and a method of measuring a coverage by precipitates of grain boundaries. Figure 3 is a graph showing a condition of the pre-weld heat treatment to the welding field destination for the purpose M ₂₃ C ₆ type carbide solid solution, the effects on the existence form of carbides where the conditions on. FIG. 4 is a graph showing the relationship between the time and heat treatment before welding for the purpose of solid solution of carbide. FIG. 5 is a graph showing the relationship between the temperature of heat treatment before welding for the purpose of solid solution with carbide and the toughness of the welded joint and the heat affected zone. FIG. 6 is a schematic perspective view of the joint for explaining the welding pre-welding butting state and part name of the welded joint, and the application range of the heat treatment before welding. FIG. 7 is a graph showing the relationship between the creep rupture strength and the heat treatment temperature after welding at a temperature of 700 ° C. for 1000 hours in a creep test of a weld joint test body. FIG. 8 is a graph showing the relationship between the average grain size (average value of equivalent circle diameters) of M ₂₃ C ₆ type carbides and TaC on the large angle grain boundaries in the weld heat affected zone and the heat treatment temperature after welding. Figure 9 is a graph showing an average particle size of the M ₂₃ C ₆ type carbide and TaC precipitated in high angle grain boundaries on the relationship of the particle surface distance of M ₂₃ C ₆ type carbide and TaC. FIG. 10 is a graph showing the relationship between the coverage of the large angle grain boundary with M ₂₃ C ₆ type carbide and TaC on the large angle grain boundary in the weld heat affected zone and the heat treatment temperature after welding. FIG. 11 is a graph showing the relationship between the roll reaction force of the hot rolling tester at the time of hot rolling, the surface temperature of the steel sheet at the time of rolling, and the state of the steel sheet in a laboratory.

First, the examination results of the present inventors and the new findings obtained from the examination results that led to the present invention will be described.
The heat resistant ferritic steel structure of the present invention is composed of a base material, a heat affected zone and a weld metal, and the shape is not particularly limited, and may be tubular or plate-like. . In addition, when the shape of the structure is tubular, the length is preferably 200 mm or more, and when it is plate-shaped, the length or width is preferably 200 mm or more.

  A necessary technique for the prevention of Type IV damage, which is the subject of the present invention, is to completely prevent the coarsening due to the HAZ thermal cycle through partial solution of carbide which is the root cause of Type IV damage. Therefore, in the present invention, in order to make the structure itself a chemical component in which Type IV damage is not likely to occur, measures for component design are taken as much as possible, and heat treatment is performed on the vicinity of the groove of the weld joint before welding. Apply and apply the conditions of the treatment concerned.

Hereinafter, the technology forming the basis of the present invention will be described together with the examination results (experimental results) of the present inventors and new findings obtained from the examination results. In addition, the experimental result shown below is obtained using the test piece produced as follows.
First, in a laboratory, a steel material having a chemical component (mass%) shown in Table 1 was melted and cast in a high frequency induction heating vacuum melting furnace having a steel material capacity of 300 kg to form a 300 kg weight steel ingot. Thereafter, the steel ingot was reheated to 1220 ° C. in an electric furnace in air atmosphere and then held in the furnace for 60 minutes, and then hot rolled to a steel plate test piece of 30 mm thickness by a hot rolling test apparatus. The hot rolling was completed at a temperature of 900 ° C. or higher and then allowed to cool. The obtained steel plate test piece is then tempered at 770 ° C. for 2 hours, and at this stage, it has a lath martensite structure, that M ₂₃ C ₆ type carbide and carbide (TaC) mainly composed of Ta are precipitated. Optical microscope, transmission electron microscope, scanning electron microscope, and electrolytic extraction residue quantitative analysis. The types of precipitates were determined by EDX analyzer (energy dispersive X-ray analysis) attached to a transmission electron microscope and pattern analysis of reflection peaks by X-ray diffraction of electrolytic extraction residue.
Here, M in “M ₂₃ C ₆ type carbide” is a total of 70% or more in total of Cr, Fe and W.

The grain boundary coverage by M ₂₃ C ₆ carbides and TaC is determined by the length occupancy of precipitates on large-angle grain boundaries by a scanning electron microscope observation image at a magnification of 10,000 times and a transmission electron microscope image of a thin film. Were determined. That is, it was assumed that the two-dimensional observation result of the grain boundary coverage was almost equal to the area occupancy of the three-dimensional grain boundary (boundary) plane, and the measured value was used as it is.

The cross-sectional schematic of the large angle grain boundary vicinity is shown in FIG. 2 for demonstrating the structure model of a welding heat affected zone, and the measuring method of the coverage by the precipitate of a grain boundary. Incidentally, _L 1 ~L ₉ of FIG. 2, the length of each precipitate (precipitate A through C), La and Lb represents a large angle grain boundary length.
The "grain boundary coverage" is, as shown in FIG. 2, the sum of precipitate lengths on the large angle grain boundary 7 (the sum of L _{1 to} L _{9 in} FIG. 2) as the sum of large angle grain boundaries. It is a value divided by (La + Lb), 100% when completely covered, and 0% when it is not covered at all.
The “grain boundary coverage” is an electron dispersive X-ray spectroscopy (EDX) analysis of particles deposited on large angle grain boundaries by a 10,000 × electron microscopic observation or a thin film transmission electron microscope analysis of 10,000 × The precipitate which can be judged as M ₂₃ C ₆ type carbide or TaC by transmission electron diffraction pattern analysis in Then, the length of the particles covering the large angle grain boundary is measured, and the measurement is performed by collecting at least 5 fields of view per sample, 5 or more test pieces per alloy, and a total of 25 samples or more in total It can be determined by observation or analysis of an electron micrograph.

  Further, the inter-particle surface distance on the large angle grain boundary is measured on the image in the manner shown in FIG. 2 from the observation result of the precipitate obtained by the above-mentioned microscopic observation, and the inter-particle distance It assumed that square distribution was carried out on the boundary surface, and it approximated easily as following formula (2).

λ _ave = 1.4 [l _s ] _ave − [d _s ] _ave (2)

The equation (2) is an empirical equation obtained by actually measuring the particle spacing on the grain boundary line and the particle spacing on the grain interface (both are the distance between the surfaces) in the present invention. Here, [l _s ] _ave is the mean particle center-to-center distance (nm), and [d _s ] _ave is the mean particle diameter (nm). λ _ave is the average particle surface distance (nm).

  Moreover, when welding, the test piece pair processed into the welding groove which made the piece width | variety of the said 30-mm-thick steel plate test piece obtained 200 mm was produced by machining. The bevel angle was 22.5 ° on one side, and a V bevel having 45 ° was formed as a bevel pair. The butt root (hereinafter referred to as the center of the sheet width direction including the point where the root was butt and the groove center) is 1 mm, the heat input is about 1 kJ / mm, and the welding speed is about 10 cm / min Then, 30 to 35 passes were raised to form a welded joint. A plurality of grooves with a total length of 400 mm were prepared, their joint characteristics were evaluated, and the structure observation of the weld heat affected zone was analyzed. The creep test was evaluated using creep test pieces having a diameter of 6 mm in parallel, a length of 30 mm in parallel, and a total length of 70 to 86 mm. Incoloy 800 (registered trademark) described in Table 2 which is a commercially available Ni-based alloy is applied to the weld metal, and an overmatch joint is formed so as not to cause breakage from the weld metal, and the HAZ characteristic evaluation is surely performed. I was able to do it.

(Heat treatment before welding)
In the present invention, for the purpose of preventing coarsening of carbides due to partial solid solution of carbides, prior to welding, welding is performed on carbides deposited on a portion (a portion corresponding to HAZ portion) to be a weld heat affected zone of a welded joint. Immediately after the heat treatment for Type IV damage prevention, the solid solution is completely resolubilized to suppress the solid solution of carbide and the precipitation reaction itself.
The heat treatment before welding aimed at this purpose is to hold the groove at a temperature of 1000 to 1100 ° C., so that the center of the plate thickness becomes a target temperature for 30 minutes or more. Although the holding time is a function of the plate thickness, it is difficult to formulate the weld since the weld does not always result in the joining of flat plates. Therefore, a thermocouple is embedded in a steel plate of the same shape (the chemical composition does not have to be within the scope of the present invention) in advance at a portion corresponding to the central portion of the plate thickness, and the entire member is heated. The temperature pattern of the heating device may be determined so as to achieve the target temperature.
In the following, the heat treatment for preventing Type IV damage, which is performed immediately before this welding, is also referred to as "pre-welding heat treatment".

FIG. 3 is a view showing the influence of the highest heating temperature holding time and the highest heating temperature holding temperature in heat treatment before welding on the form of carbides. In this test, a plurality of partially cut steel plate test pieces of 50 mm square and 10 mm thickness are prepared, heat-treated only for various temperatures and times, and then cut and processed for cross section, and the presence or absence of precipitation of carbides by transmission electron microscopic structure observation confirmed. “●” in the figure is an example in which M ₂₃ C ₆ type carbide and TaC remained undissolved in the middle of decomposition in the specimen after heat treatment and cooling after being allowed to cool, “○” is solid solution of all carbides The case was not observed at all. “□” indicates M ₂₃ C ₆ type carbides, but it is an example where stable M ₂₃ C ₆ type carbides having the same form as the base metal and not undergoing decomposition solid solution were found (however, precipitation amount In this example, it is only 0.3 pieces / μm ² , and the deformation resistance of the creep strength generated by this carbide is extremely small at 8 to 10 MPa. Confirmed by theoretical calculation of threshold stress value from particle density).

Separately, when a creep test was carried out in a welded joint subjected to the above-mentioned heat treatment, among the forms of the above-mentioned carbides, the precipitate form which induces Type IV damage is not dissolved during the decomposition shown by “●”. It was revealed that the remaining carbide alone was not generated in the other form or in the case where no carbide was observed. That is, in the heat treatment before welding, the M ₂₃ C ₆ type which remains undissolved in the middle of decomposition by holding at least 30 minutes at a temperature of 1000 ° C. or more for at least 30 minutes and longer than necessary time It was found that carbides and TaC were not observed at all, and it was possible to decompose and re-solidify all.
In the temperature range of 1000 ° C. or less, M ₂₃ C ₆ type carbide and TaC are not solid but not solid carbides, but thermodynamically stable carbides (equilibrium precipitation carbides) that do not decompose and solid solution to high temperatures. it is conceivable that. That is, according to the experiment, it was found that the M ₂₃ C ₆ type carbide and TaC precipitated at the time of tempering treatment of the base steel plate can not be substantially completely re-dissolved by the heat treatment which does not satisfy the necessary time. .

  Also, according to the graph shown in FIG. 3, since the non-solid solution carbide causes Type IV damage, the maximum heating temperature is kept within the temperature range of 1000 ° C. to 1100 ° C., and maintained at this maximum heating temperature for 30 minutes or more It can be seen that the heat treatment conditions of the above are essential conditions. However, when the maximum heating temperature of the heat treatment before welding exceeds 1100 ° C., the heat-resistant steel within the scope of the present invention rapidly increases the grain size of γ grains, resulting in a decrease in toughness and Toughness is a problem.

In addition, the relationship between the holding time and the plate thickness in the heat treatment before welding needs to be heated locally and held so as to be 30 minutes or more for a longer time than the plate thickness corresponding heat treatment time Ts determined by the following equation (1) There is. Only by this treatment, M ₂₃ C _6- type carbides and TaC are re-reformed at a stage prior to welding to such an extent that coarsening through non-solid solution of carbides generated during a short heat effect can be prevented which causes Type IV damage. It is possible to cause solid solution. In addition, since the following formula (1) assumes application to heat-resistant steel, the minimum board thickness assumes 3 mm.

Ts = 30 × (t / 25) ^1.2 +28... (1)

  FIG. 4 shows the relationship between the thickness t and the heat treatment temperature holding time before welding. The curve in the figure is the boundary represented by the equation (1), and when heat treatment before welding is performed, the condition is determined so that the retention time above the curve is obtained.

  FIG. 5 is a view showing the relationship between the temperature of heat treatment before welding and the toughness (Charpy absorbed energy) at 0 ° C. of the heat affected zone at welding. In this experiment, the temperature of the heat treatment before welding was variously changed as shown in the graph, and the plate thickness was 15 mm, and the heat treatment time (holding time) was unified at 60 minutes. The Charpy impact test was carried out according to the method described in JIS Z2242 by collecting a test piece of 55 mm in total length obtained by processing a 10 mm square, 45 mm 2 mm V notch in the plate length direction from the thickness center position and measuring absorbed energy at room temperature. . The notch is processed so that it is horizontal to the plate thickness direction with the point at which the interface between the weld joint metal and the weld heat affected zone intersects the line at the plate thickness center (plate thickness center position with reference to the interface). . It is clear that the heat affected zone toughness of the welded joint subjected to the pre-weld heat treatment above 1100 ° C. is lower than the target 27 J.

From the above, the heat treatment before welding in the present invention is performed in the temperature range of 1000 ° C. to 1100 ° C. for 30 minutes or more longer than the heat treatment time Ts determined by Equation (1) above. Local heating and holding is important.
Note that there is no limitation on the extension of the holding time in the heat treatment before welding, and it may be appropriately determined according to the thickness of the target steel material, but it is substantially the longest from the viewpoint of avoiding damage due to high temperature oxidation of the groove. 5 hours is the limit.

From the examination results of the present inventors described above, by performing the heat treatment before welding under the above conditions, the carbides at the stage where the welding heat cycle is applied are decomposed by the short time heating at the fine grain area equivalent heat cycle temperature. The solid solution does not exist in advance in the groove of the welded joint, and naturally, there is no possibility that coarsening of undissolved carbides will occur in the subsequent post-welding heat treatment. Incidentally, precipitation of the M ₂₃ C ₆ type carbides so that the resulting first time in the heat treatment after welding.
Therefore, in the heat-resistant steel of the present invention range in which B is added by 0.008% or more, fine grain structure is not generated, and therefore, in HAZ, coarse structure due to thermal cycle of carbide is avoided while having the same structure structure as the base material. Precipitation state is obtained.

Next, a method of applying heat treatment before welding will be described.
In the present invention, it is important to apply the heat treatment before welding only to the vicinity of the welding groove. FIG. 6 is a schematic perspective view of the joint for explaining the butting condition before welding and the names of each part of the welded joint, and the heat treatment application range before and after welding.
Specifically, as shown in FIG. 6, the heat treatment before welding is applied only to a portion within 20 mm to 100 mm on one side from the welding groove surface 12.

  As for the method of applying heat treatment, since the heat treatment is local heating and intended to improve the on-site workability, resistance heating bands such as “high frequency induction heating”, “electric current heating”, and a band heater (open It can be carried out by direct winding or coating (limited to those protected by an insulator which is not electrically conductive). All these heating methods are selected from the viewpoint of having a function capable of easily performing the groove heating at the boiler construction site, but the method of applying the heat treatment is appropriately selected within the range that does not impair the effects of the present invention. You may

Further, the heating range of the heat treatment is limited to a range of 20 mm to 100 mm on one side from the weld groove surface 12 and 50% or less of the length or width of the structure. The reason for this range is that the width (HAZ width) outer edge of the weld heat-affected zone from the center of the weld groove to the base material is only a few tens of mm when using ordinary arc welding and exceeds 100 mm It is because there is substantially no thing. Further, since the high-strength ferritic heat-resistant steel structure of the present invention uses only a welding heat input of up to several tens of thousands of kJ / mm when welding, the weld metal is substantially even if the HAZ width is an extremely thick member Most of the cases are 30 mm or less including the width. Therefore, the heating range of the heat treatment is preferably further reduced according to the output of the heating apparatus and the ease of handling, in view of the process, and it is further preferable to set it to 80 mm or less. Still more preferably, it is 50 mm or less.
When heat treatment before welding is performed, residual γ is generated. Since the residual γ induces deformation and cracking of the member, the smaller the amount of generation, the better. From this point of view, the application range of the heat treatment before welding is limited to 50% or less of the length or width of the structure.

When the width occupied by the weld metal is large or when the welding heat input is large, the position of the weld heat affected zone width is located at a position where the position of the weld heat affected zone width exceeds 20 mm on one side from the groove center There are times when That is, when only the range of less than 20 mm is heated from the position (reference numeral 13) where the weld groove surface 12 and the surface of the base material contact (intersect) as the heating range of heat treatment before welding, the heat treatment before welding is the welding heat affected zone It may not reach the area corresponding to the fine grained area, and as a result, Type IV damage may occur. Therefore, the lower limit value of the heating range of the heat treatment before welding is limited to within 20 mm from the welding groove surface 12. The effect of the present invention can be achieved even if the heating range of the heat treatment before welding exceeds 100 mm from the position (reference numeral 13) where the weld groove surface 12 and the surface of the base material contact (intersect) If it is necessary to heat-treat a wide width by a simple heat-treatment apparatus, and if it is a steel pipe or steel member substantially having a diameter slightly larger than 200 mm, the same effect as total heating and local heating described later The heating range of the heat treatment before welding is set to 100 mm or less from the weld groove surface 13 from the viewpoint of the implementation cost, such as the effect of the heat treatment time being not substantially obtained and the heat treatment time becoming long.
In addition, the code | symbol 14 in FIG. 6 shows the minimum width which should give the heat processing before welding, and the code | symbol 15 shows the maximum width which should give the said heat processing.

Above-described as, the heat treatment conditions and the heat treatment method for imparting coarsening of M ₂₃ C ₆ type carbide, which causes, Type IV damage can be sufficiently prevented state.

(Post-weld heat treatment and metallographic structure)
Next, the metallographic structure and heat treatment after welding of the high strength ferritic heat resistant steel structure according to the present invention will be described.
The post-weld heat treatment of the present invention refers to a portion of 20 mm or more and 100 mm or less and 50% or less of the length or width of the structure from 720 ° C. or more to 760 ° C. It is a process which heats below and hold | maintains to this temperature for 1 hour or more.

Post-weld heat treatment (hereinafter also referred to as PWHT) is generally applied at a temperature of (tempering temperature of base material-20) or less and for a time corresponding to the plate thickness, but it is the same as the base material In order to exhibit creep rupture strength, it is necessary to control the precipitation state of M ₂₃ C ₆ type carbide and TaC to the same level as the base material. Preferably, the heat treatment temperature after welding is equal to the tempering temperature of the base material before welding, but the heat treatment after welding simultaneously heats the weld metal.

  The weld metal is usually subjected to a post heat treatment temperature (tempering temperature of the base material of −20 ° C. so that the as-solidified structure obtained by welding can improve the creep strength of the joint, in particular to avoid a decrease in the high temperature strength of the weld metal. ) The alloy design is based on the following post-welding heat treatment temperature. That is, the alloy of the weld metal selects a chemical component that tends to have a higher high temperature strength than the base material. However, since the chemical composition of the base material is optimized to improve the creep rupture strength after 100,000 hours, the weld metal which emphasizes the initial high temperature strength affects the allowable stress of the standard as a whole joint. It is general that the creep strength is low in the non-range (having a value within the average value of breaking strength-5%). That is, carrying out PWHT at high temperature promotes the softening of the weld metal and is likely to cause a decrease in the creep rupture strength of the weld metal. Therefore, it is not preferable from both of the viewpoints of creep strength as a joint to apply the same post-welding heat treatment as the tempering condition of the base material. In addition, even when the tempering temperature is too low, the precipitation state of HAZ is not equal to that of the base material, and the creep strength is lowered.

Therefore, in the present invention, the temperature of the heat treatment after welding is strictly limited, and the metal structure, in particular, the precipitates are controlled to solve this problem.
The present inventors examined the precipitation morphology of M ₂₃ C ₆ type carbides and TaC in the weld heat affected zone in order to improve the creep rupture strength of the welded joint. As a result, the average equivalent circular particle diameter of M ₂₃ C ₆ type carbide and TaC is 300 nm or less, the coverage of large angle grain boundary with M ₂₃ C ₆ type carbide and TaC is 40% or more, and large angle grain It has been found that the inter-surface distance of M ₂₃ C _6- based carbides and TaC precipitated on the boundary is required to be 150 nm or less in HAZ after heat treatment after welding regardless of the particle type.
In the present invention, the "large angle grain boundary" refers to a crystal boundary or grain boundary in which adjacent crystal orientations are relatively 15 ° or more as compared around the <110> common rotation axis.

  That is, in the present invention, the heat treatment after welding is held at a temperature range of 760 ° C. or less and 720 ° C. or more for one hour or more. Thereby, the area ratio of retained austenite can be made 0.2% or less in a portion of 50% or less of the length or width of the heat resistant steel structure, and the precipitate can be made into the desired form described above.

Figure 7 shows the relationship between the heat treatment temperature after welding and the 700 ° C, 1000 hour creep rupture strength (650 ° C, simplified temperature accelerated creep test equivalent to a 100,000 hour creep test) of a welded joint having components within the scope of the present invention. Show. The holding time at the heat treatment temperature after welding was one hour in the graph shown in FIG.
As shown in the graph of FIG. 7, when the heat treatment temperature after welding (PWHT temperature) exceeds 760 ° C., the creep rupture strength of the welded joint is reduced, and in particular, it is understood that the weld metal fractures (● in the figure) frequently occur ( PWHT temperature too high)

FIG. 8 shows the relationship between the PWHT temperature and the average value of the equivalent circular particle sizes of M ₂₃ C ₆ type carbides and TaC in the heat affected zone of the welded joint (hereinafter also referred to as the average particle sizes of M ₂₃ C ₆ type carbides and TaC) Is a graph showing Although the holding time at the PWHT temperature was 1 hour to 4 hours in the graph shown in FIG. 8, the change in the holding time hardly affected the average grain size of M ₂₃ C ₆ type carbide and TaC, so We do not organize every holding time in the inside.
As apparent from FIG. 8, when the PWHT temperature exceeds 760 ° C., the average particle size of the combined M ₂₃ C ₆ type carbide and TaC becomes coarser than 300 nm, the creep rupture strength decreases, and the Type IV damage becomes It turns out that it occurs. In the heat treatment below the transformation point, the crystal structure, particularly the size of the large-angle grains, hardly changes, and at the same time, the coverage by precipitates on grain boundaries is also affected.

FIG. 9 is a graph showing the relationship between the inter-surface distance between particles of M ₂₃ C ₆ type carbide or TaC on the large angle grain boundary and the average particle diameter. In addition, FIG. 9 is the result of having considered PWHT temperature as 740 degreeC, and holding time being 1 hour.
As apparent from FIG. 9, regardless of the type of M ₂₃ C ₆ type carbide and TaC, it is understood that the distance between particle surfaces is 150 nm or less when the average particle diameter is 300 nm or less. That is, when the heat treatment temperature after welding is 760 ° C. or less, the average particle diameter of the M ₂₃ C ₆ type carbide and TaC is 300 nm or less, and the particle surface distance on the large angle grain boundary is 150 nm or less I understand. However, when the average particle size exceeds 300 nm, the distance between particle surfaces exceeds 150 nm, and at this time, reinforcement by intergranular precipitates of the welded joint can not be expected during creep deformation in a long time, and Type IV damage occurs Do.

The creep rupture strength also decreases when the PWHT temperature is 720 ° C. or less. The reason is that tempering of the heat affected zone is insufficient, and the above-mentioned large angle grain boundary occupancy of M ₂₃ C ₆ type carbide and TaC described below does not exceed 40% (ie, Type IV damaged example ).
FIG. 10 is a graph showing the relationship between the grain boundary coverage by the precipitates (M ₂₃ C ₆ type carbide and TaC) deposited on the large angle grain boundaries in the heat affected zone of the welded joint and the PWHT temperature. The holding time at the PWHT temperature was one hour in the graph shown in FIG.
As apparent from FIG. 10, when the PWHT temperature is less than 720 ° C., the grain boundary coverage is clearly less than 40%. In this case, Type IV damage inevitably occurs in the joint. On the other hand, when the PWHT temperature is 760 ° C. or higher, the grain boundary coverage decreases rather than 40% due to accelerated coarsening of M ₂₃ C ₆ type carbide and TaC, and Type IV damage occurs. I also understand.

As for the time of PWHT, M ₂₃ C ₆ type carbide and TaC are precipitated at grain boundaries, and the equivalent circle diameter is 300 nm or less according to electron microscopic observation, and the coverage on the large angle grain boundary (boundary) is 40% or more. However, since the heat treatment has the purpose of relieving residual stress of the joint, it is at least 1 hour or more.

Generally, the construction conditions may be determined by the function of the thickness t of the steel plate, and the required PWHT time (TP) is
TP ((t / 25 + 1) x 60 minutes ... (c)
The outline of the
The present invention can also roughly follow this formula. However, it is preferable to set the heat treatment for a maximum of 168 hours (one week) because heat treatment for too long time causes an increase in construction cost.

Here, the equivalent circle particle size of the precipitate on the large angle grain boundary was determined as follows.
First, the cross-sectional structure of the test piece after heat treatment after welding is observed with a scanning electron microscope, and then the grain boundary structure constituting the ferrite structure is made more detailed in EBSP (electron beam backscattering diffraction pattern analyzer) I observed it. Among them, a block grain boundary specific diffraction angle, that is, the angle around the <110> common rotation axis between the adjacent crystal orientations and having an adjacent crystal orientation difference of 15 ° or more, ie 54 Image processing of °, 60 °, and 16 ° is determined as "block grain boundary (large angle grain boundary)", and the crystal structure of M ₂₃ C ₆ type carbide and TaC precipitated on the large angle grain boundary Only the precipitates having the same electron diffraction pattern were identified. Then, based on an electron micrograph of 10000 times of the precipitates, the diameter of particles on the cross section was determined on the photograph.
Image picture of 10000 times observed five or more views in the heat affected zone of one joint, measure the cross-sectional area of the precipitate for all the particles, and assume that this is all a circle. Equivalent diameter.

The greatest feature of the present invention is the combination of the aforementioned "pre-welding heat treatment for the purpose of completely preventing the coarsening of carbides in HAZ by welding heat cycle" and the present chemical components which exert the heat treatment effect. Among the most important pre-welding heat treatment applications, the greatest difference from the techniques described in Patent Documents 1 to 4 lies in the limitation to “local heating”. Although the effect of the application of this "local heating" has already been described, it is another feature that "local heating" is more advantageous in creep strength than "total heating of the welding member".
The reasons why "local heating" is advantageous in creep strength are discussed in more detail below.

The techniques described in Patent Documents 1 to 4 perform “heat treatment before welding” on the entire welding member, for example, a steel pipe (12 m in the case of a heat exchanger) after groove processing of welding. At this time, residual γ is generated in all the steel pipes.
With such a large-sized member, the material nonuniformity in the steel material manufactured by the steelmaking consistent process can not be completely wiped off. In the case where the forging process is applied after the electric furnace steelmaking method, it is more remarkable. For example, elements such as W and V added in the present invention have segregated portions where the concentration is high.
That is, after forming a residual γ by heating the whole of the member having such a segregated portion, it is possible to prevent the formation of fine grained area in the welded joint and to prevent the coarsening of M ₂₃ C ₆ type carbide and TaC. Even if the subsequent PWHT is not carried out for the entire part after welding, the base metal which is not affected by the welding heat has a high B concentration in the base metal, so the low temperature short of heat treatment after welding This is because the residual heat treatment can not be completely decomposed by the heat treatment for a long time, resulting in a structure in which the residual light remains.

If it is desired to completely decompose the residual γ of the entire steel material, in a high Cr steel of 9% or more, heat treatment must be performed at 760 ° C. or more for at least 2 hours to cause sufficient diffusion of segregated elements to occur. It became clear by our studies. However, as described above, when PWHT is performed on the entire steel material at 760 ° C., which is lower than the tempering temperature of the base material, complete decomposition of the residual γ may not be realized.
That is, in the case of "total heating of the welding member", the heat treatment after welding is also total heating, and in order to avoid the residual γ remaining in the base material, the extremely limited high tempering temperature and PWHT temperature are selected. There is a need. However, the actual heating capacity of the PWHT installation apparatus is that the temperature accuracy is ± 20 ° C., and it does not necessarily correspond to strict temperature control.
From this, when heat treatment before welding is performed on the entire member, the possibility of the residual γ of the base material can not be completely denied in consideration of the on-site workability. That is, when such a steel material is applied to a high-temperature device (for example, a boiler or the like) with a considerable amount of residual γ present, deformation or cracking due to the residual γ may occur.

  Since the working temperature of high-temperature pressure equipment such as a boiler is about 600 ° C. at present, decomposition of residual γ takes time, and in some cases, it is considered to require several thousand hours. During this time, residual γ is gradually transformed into ferrite and carbide in the use environment to cause thermal expansion. Since this local thermal expansion naturally induces deformation or cracking of the member, the presence of residual γ is not permitted. That is, materials that may have residual γ may not be applicable to the real environment.

  On the other hand, in the case of "local heating", there is no residual γ formation of the base material. Moreover, since only a portion within 100 mm from the grooved surface of the weld joint is heated, there is a possibility of generation of residual γ at the portion, but temperature control is extremely easy because the heating range is narrow, and overall heating The probability of existence of residual γ is much smaller than in the case of. In addition, even if there is a slight residual γ, since it exists only within 100 mm from one side of the groove surface, the effect of thermal expansion is ensured by securing the base metal portion without residual γ Can be significantly reduced.

This is noticeable in so-called "round weld joints" where the weld line is perpendicular, for example, to the tube axis.
On the other hand, when the weld line is parallel to the tube axis, it can be seen in the piping of the boiler for power generation, etc., but this is a large diameter steel pipe, and if there is a diameter of 350 to 700 mm, the circumferential length will be 1100 to 2200 mm. Even if thermal expansion occurs in a portion of 100 mm on one side and 200 mm at the maximum among them, as described above, it is easy to relieve the expansion stress in another portion (for example, base material). When the heat treatment width before welding is narrowed in consideration of the actual welding heat affected zone width, the influence can be further reduced.
However, when attempting to locally heat a wide region exceeding 100 mm on one side, a large amount of power is required in the simplified heat treatment such as high frequency induction heating or a band heater assumed by the present invention, which causes a cost problem. In addition, reheating of a wide range is disadvantageous in that the time of the heating process is prolonged. Furthermore, when the size of the actual member is about 200 mm, for example, if a steel pipe having a diameter of 200 mm is heated, a post-welding heat treatment, which agrees with the entire heating, is eventually performed. As a result, the simplicity of the present invention is lost, and when residual γ occurs outside the heat affected zone, there is a concern that expansion or deformation due to the progress of transformation during use or destruction due to thermal stress may occur. Become.

  In the present invention, the influence of residual γ formed on the base material when the ratio of the construction width (length) to the member including the welded joint of heat treatment before welding exceeds 50% of the width or length of the structure. Even when the amount of residual γ produced is as small as 0.2% by volume, the creep life is shortened due to the thermal stress cracking caused by the decomposition of the residual γ under creep conditions, and the creep rupture strength is consequently reduced. It will be a thing. In other words, if the area ratio of the residual γ in the metal structure is 0.2% or less at 50% or more of the volume of the structure, cracking can be prevented due to the thermal expansion caused by the residual γ. Needless to say, the wider the region where the area ratio of the residual γ is 0.2% or less, the better, and it is preferably 100% of the volume of the structure.

  On the other hand, in general, heat treatment by "local heating" always has the problem of a common "intermediate temperature heating zone (intermediate temperature range)". That is, when only the target part is heated to the target temperature, an intermediate temperature range heated to a temperature zone lower than the target temperature is generated in the part adjacent to the target area, and the target tissue and effect are It is always necessary to take into account the possibility of non-emerging sites and the possibility of producing special tissue changes at cold sites.

However, in the case of the present invention, another feature is that this problem does not occur.
That is, when the heat treatment before welding is performed at 1000 to 1100 ° C., the adjacent portion is reheated to a temperature range of 1000 ° C. or less. In such a case, the possibility of incomplete solid solution of carbide may also be considered. However, incomplete solution of carbides occurs only when heating is performed for only a few seconds, such as the effect of welding heat, that is, when the time necessary for solution and dissolution of carbides is not sufficiently given. If, as long as phase transformation occurs, sufficient time for decomposition and solid solution is given as in the present invention, the solid solution limit of carbon in the γ phase is much larger than that in the α phase, so The dissolution and dissolution of C. is completed, and incomplete solution carbide does not occur. Therefore, the heat treatment before welding according to the present invention is characterized in that there is no concern of coarsening of carbides in the intermediate temperature range.

  It should be noted that in the case of conventional steels, it may be sufficient that the refinement of the structure occurs in the intermediate temperature range, but the heat-resistant steel according to the present invention has a high B concentration and does not produce fine grain areas. It is theoretically clear that the grain structure becomes equivalent to that of the base material also by not grain refining due to the occurrence of γ transformation.

With regard to performing post-weld heat treatment locally, it is not necessary to consider the intermediate temperature range for the purpose of stress relief annealing, and M ₂₃ C ₆ type carbide and TaC are precipitated at grain boundaries of a structure similar to that of the base material. It is also apparent that there is no problem at all if the temperature after heat treatment after welding is strictly adhered to, since heat treatment is carried out as far as the heat treatment is carried out including the base metal portion which is not affected by the welding heat.

  That is, in the chemical composition of the present invention and heat treatment before welding, it is not necessary to consider heating in the intermediate temperature range, it is effective to carry out heat treatment before welding by "local heating", and "local heating" This method is desirable in that it can reduce the influence of residual γ on "overall heating", and is advantageous in cost and comprehensively excellent.

The structure according to the present invention is characterized in that the applied width or length of the heat treatment before welding is 50% or less of the width or length of the structure, but the portion subjected to the heat treatment before welding is the following It can be identified in the same way.
In the case where the “execution width or length of heat treatment after welding” is smaller than the “width or length of heat treatment before welding”, the obtained structure has a region where M ₂₃ C ₆ type carbides are not precipitated It will be done. Therefore, when there is a region in which the M ₂₃ C ₆ type carbide is not precipitated, it can be determined that the region is a boundary as to whether or not the heat treatment before welding is applied.
On the other hand, the grain boundary coverage by M ₂₃ C ₆ type carbide can not be improved to 40% or more even if heat treatment after welding is applied to a portion not subjected to heat treatment before welding. Therefore, in the case where the “execution width or length of heat treatment after welding” is larger than the “width or length of heat treatment before welding”, the boundary indicating whether or not the heat treatment before welding is performed is M ₂₃ C ₆ type carbide It can be judged by whether or not the grain boundary coverage by B is 40% or more.
Incidentally, the grain boundary coverage by the presence or M ₂₃ C ₆ type carbide of the M ₂₃ C ₆ type carbide, as described above, can be confirmed by transmission electron diffraction pattern analysis in electron microscopy or thin film transmission electron microscopic analysis.

  Next, the chemical components of the heat resistant steel of the present invention will be described. In addition, "%" of content of each element means "mass%."

C: 0.01 to 0.12%
C is an element that forms carbides and enhances the hardenability. In the present invention, in order to improve creep rupture strength, the amount of C is made 0.01% or more. In order to enhance the precipitation strengthening ability, it is preferable to add 0.05% or more of C. On the other hand, if the amount of C is too large, the precipitates become coarse and the occupation ratio of grain boundaries decreases, so the amount of C is made 0.12% or less. Further, if the amount of C is excessive, carbides formed in grain boundaries may be coarsened, which may lower the creep rupture strength. Therefore, the amount of C is preferably made 0.10% or less.

N: 0.003 to 0.015%
N is an element that forms a nitride, and is an element effective to precipitate VN and improve the initial creep strength. In order to receive this effect, N is contained 0.003% or more. In addition, Al mixed from a refractory or the like may combine with N, and the amount of N for producing VN may not be sufficiently secured. In consideration of such a case, the amount of N is preferably 0.005% or more. However, if the N content exceeds 0.015%, BN may precipitate, so the upper limit is made 0.015%. Further, since N is an element which is activated by neutron irradiation to embrittle the steel, when the heat resistant steel is used in a nuclear power plant etc., the amount of N may be made 0.010% or less. preferable.

B: 0.008 to 0.050%
B is an element that enhances the hardenability of the steel material in a solid solution state (including a state of being segregated at grain boundaries) and forms a martensitic structure and a lower bainite structure having a high dislocation density. In the present invention, B dissolves in carbides and intermetallic compounds to increase the thermal stability, delays the coarsening of these precipitates, or precipitates as a boride to increase the precipitation strengthening ability. Element. In order to improve creep rupture strength, B needs to be added at 0.008% or more, and addition of 0.012% or more is preferable. On the other hand, if B is added excessively, the weldability deteriorates, so the B content needs to be 0.050% or less. When it is necessary to increase welding heat input, it is preferable to make B amount into 0.017% or less. A more preferable B amount is 0.015% or less.

V: 0.10 to 0.50%
V is an element that combines with N to form a nitride, and composite precipitates in the grains in alignment with NbC. In the present invention, V of 0.10% or more is added to enhance creep rupture strength. In order to enhance the effect of precipitation strengthening, the addition of 0.15% or more of V is preferable, and the addition of 0.17% or more of V is more preferable. On the other hand, when V exceeding 0.50% is added, coarse VCs precipitate to affect the toughness, so the V amount is made 0.50% or less. In order to enhance the toughness, the V amount is preferably 0.40% or less, and more preferably 0.35% or less.

Si: 0.02 to 0.45%
Si is a deoxidizing element, and 0.02% or more is added. In order to enhance the effect of deoxidation, it is preferable to add 0.10% or more of Si. Moreover, Si is also effective in improving the oxidation resistance, and it is more preferable to add 0.20% or more. On the other hand, when Si is added in excess of 0.45%, an oxide containing Si may become a starting point of brittle fracture and impair toughness. In addition, excessive addition of Si substitutes for the site of solid solution W to promote the precipitation of Fe ₂ W, which may lower the creep rupture strength. Do. In order to enhance the toughness, the amount of Si is preferably 0.40% or less, more preferably 0.35% or less.

Mn: 0.20 to 0.60%
Mn is a deoxidizer, and 0.20% or more is added in the present invention. It is preferable to add 0.35% or more of Mn because the toughness decreases if the deoxidation is insufficient. On the other hand, Mn is an austenite-forming element, and the mobility of dislocations is increased to locally accelerate the recovery of the structure. Therefore, if it is added excessively, the creep characteristics deteriorate. In the present invention, in order to secure creep strength, Mn is made 0.60% or less. In order to further increase the creep rupture strength, the Mn content is preferably 0.55% or less, more preferably less than 0.50%.

Cr: 8.0 to 12.0%
Cr is an important element that enhances the hardenability of steel materials and causes precipitation strengthening of steel materials as carbides. In order to obtain high creep rupture strength at high temperatures of 650 ° C. or higher, it is necessary to secure the amount of M ₂₃ C ₆ type carbide mainly composed of Cr, promote coarsening, and suppress excessive coarsening. In the present invention, 8.0% or more is added. In consideration of the anti-steam oxidation property, it is preferable to add 8.5% or more of Cr. On the other hand, excessive addition of Cr accelerates coarsening of the M ₂₃ C ₆ type carbide at a temperature of 650 ° C. and creep characteristics deteriorate, so the amount of Cr is made 12.0% or less. The amount of Cr is preferably 10.5% or less, more preferably 9.20% or less.

W: 2.00 to 7.50%
W is an element that forms an intermetallic compound with Fe and contributes to the improvement of creep characteristics. When 2.00% or more of W is added, intermetallic compounds are precipitated during long-term use, which greatly contributes to creep rupture strength. Moreover, in order to improve the precipitation density to the grain boundary, addition of 2.5% or more of W is preferable, and addition of 2.7% or more of W is more preferable. On the other hand, when W is added excessively, the coarsening of the Fe ₂ W type intermetallic compound (Laves phase) becomes faster, so the W amount is made not more than 7.50%. In order to suppress the coarsening of the Laves phase, the W content is preferably 7.0% or less, more preferably 6.5% or less.

Nb: 0.02 to 0.10%
Nb is an element that forms a carbide, and precipitates in the grains to contribute to the improvement of the creep rupture strength. If NbC type carbide and NN complex precipitate with VN, the movement of dislocation can be effectively suppressed, so in order to maintain the creep strength in a relatively short time, the Nb content is made 0.02% or more. Moreover, in order to improve the intragranular precipitation strengthening ability by NbC, addition of 0.03% or more of Nb is preferable, and addition of 0.04% or more of Nb is more preferable. On the other hand, when Nb is added in excess of 0.10%, it precipitates as coarse NbC and loses toughness, so the Nb amount is made 0.10% or less. In order to precipitate NbC finely, it is preferable to make the amount of Nb 0.08% or less, more preferably 0.07% or less.

Ta: 0.10 to 1.00%
Ta, like Nb, is an element which forms carbides. Since TaC is thermodynamically unstable than NbC, it does not precipitate at a very high temperature, but mainly precipitates in the ferrite region, that is, at a temperature below the transformation point. For this reason, the precipitation based on the body diffusion limited control of Ta becomes the main, the precipitation position differs from NbC, and it becomes both the grain boundary and the inside of the grain. Among them, TaC deposited at grain boundaries has the effect of discontinuously depositing together with the Laves phase to increase the grain boundary coverage, similarly to M ₂₃ C ₆ type carbides. That is, it is an additive element that is the core of the present invention. Intergranular precipitation of TaC improves the resistance to Type IV damage of steels within the scope of the present invention over conventional steels as long as the construction procedure in conventional weld joints is taken, but it is possible to completely suppress this. Is not a stable precipitate and coarsens like other carbides in long-term creep test.
In order to utilize Ta as carbides for precipitation strengthening, 0.10% of addition is required, and preferably 0.15% is added. In addition, when it is desired to exhibit long-term creep rupture strength which makes the best use of grain boundary precipitation strengthening and improve characteristics such as toughness by another technique, addition of 0.20% or more is preferable.
In addition, a large amount of addition of Ta causes a marked decrease in hot workability. This is due to coarse precipitation of TaC at grain boundaries, and in order to avoid this, the maximum addition amount is made 1.00%. If the high temperature strength of the material is relatively small, processing cracking may occur during hot working. Therefore, when the amount of C added is relatively low, the amount of Ta is preferably 0.70% or less, more preferably 0. 50% or less.

Co: 0.50 to 3.00%
Co is an austenite stabilizing element, and is an element that improves the hardenability and enhances the toughness. Co is the only element that does not change the transformation point, and the present inventors have also found the effect of Co to reduce the dislocation mobility. In the present invention, 0.50% or more of Co is added to increase creep rupture strength and to suppress formation of a ferrite phase (δ ferrite). In order to enhance these effects, the amount of Co is preferably 0.7% or more, more preferably 1.0% or more. On the other hand, Co promotes the precipitation of the σ phase, so the amount of Co is made 3.0% or less. The amount of Co is preferably 2.5% or less, more preferably 2.3% or less. In an environment where neutrons are irradiated, such as nuclear power plants, Co may be activated and toughness may be impaired by neutron irradiation embrittlement. In such a case, the amount of Co is 1.50% or less. It is preferable to

Nd: 0.005 to 0.05%
Nd is an element that can be combined with Mn to prevent the formation of MnS that extremely adversely affects toughness, for example, as inclusions in steel. It is also an element which is easy to be accumulated in grain boundaries and which is effective in preventing grain boundary segregation and reheat embrittlement by alleviating grain boundary segregation of S. In the present invention, when Nd is accumulated at grain boundaries and the grain boundary coverage intended by the present invention is improved, precipitation of M ₂₃ C _6- type carbides and precipitation of Fe ₂ W-type Laves phase cause impurities due to impurities Cr. And Fe is not consumed, for example, in the form of sulfide or phosphide. This effect also works on Ta precipitated at grain boundaries. That is, Nd is also an important element in enhancing the effect of the present invention.
Such an effect of alleviating impurity segregation at grain boundaries is manifested by adding Nd of 0.005% or more. In the case where the concentration of S and P, which are elements that easily cause grain boundary segregation, is relatively high, it is preferable to add 0.010% or more. More preferably, 0.015% or more is added.
When Nd is added at a high concentration, NdS is formed, which may rather promote intergranular embrittlement. Therefore, the upper limit of addition is made 0.05%. In the case of a high oxygen concentration where the toughness is lowered by the formation of NdO, the addition of 0.045% or less is preferable.

  In the present invention, the content of cold iron source such as scrap and Mo, Ni, Cu and Al mixed as impurities from a refractory is limited to the following range.

Mo: 0.030 or less Mo partially dissolves in Fe ₂ W and carbide M ₂₃ C ₆ and at the same time forms carbides of the Mo ₂ C and Mo ₆ C types and promotes the coarsening of precipitates, The content is limited to 0.05% or less in order to adversely affect long-term creep characteristics. It is preferable to limit the Mo content to 0.03% or less because borides mainly composed of Mo tend to coarsen and reduce creep rupture strength. Moreover, since Mo is an element which is activated by neutron irradiation to embrittle steel, it is more preferable to limit the amount of Mo to 0.01% or less when used in a nuclear power plant or the like.

Ni: 0.10% or less Ni is an element effective for improving toughness and stabilizing austenite, but it increases the mobility of dislocations and significantly reduces the creep rupture strength. Limit the amount. In the present invention, the amount of Ni is limited to 0.10% or less in order to suppress a decrease in creep rupture strength for a long time. In order to enhance the creep characteristics, the content of Ni is more preferably limited to 0.05% or less, still more preferably to 0.03% or less.

Cu: less than 0.05% Cu is an element effective for the stabilization of austenite, but when it is produced by normalizing-tempering as in the present invention, it is used as ε-Cu (metallic Cu) in steel It precipitates alone. When heated to 1100 ° C or higher during hot working, iron is selectively oxidized, and when Cu gathers at grain boundaries, a local low melting point metal accumulation zone is formed, which induces intergranular delamination cracking. (Red-hot brittle) may occur. Thus, in the present invention, Cu has a small contribution to the stabilization of austenite, and the hot workability is reduced, so the amount of Cu is limited to less than 0.05%. In order to enhance the manufacturability, the content of Cu is more preferably limited to 0.03% or less, still more preferably 0.01% or less.

Al: 0.005% or less Al binds to N in the present invention, inhibits precipitation strengthening by VN, and reduces the effect of intragranular strengthening, so the content thereof is limited to 0.005% or less. Since the creep rupture strength is reduced by a small amount of Al, the Al content is preferably limited to 0.003% or less, more preferably 0.002% or less.

  Further, in the present invention, since P, S and O are impurities, the content is limited as follows.

P: 0.020% or less P segregates at grain boundaries, promotes intergranular fracture and lowers toughness, so the content is limited to less than 0.020%.

S: 0.010% or less S combines with Mn to limit the content to less than 0.010% in order to prevent the decrease in toughness due to the formation of coarse MnS.

O: 0.010% or less O forms an oxide cluster that is a starting point of brittle fracture and reduces the toughness, so the content is limited to less than 0.010%.

  In the present invention, if necessary, one or both of Ti and Zr can be added to fix N.

Ti: 0.005 to 0.15%
Ti is an element having a very high affinity for N compared to B. In order to suppress the precipitation of BN by the formation of TiN and enhance the effect of B to suppress the coarsening of carbides, it is preferable to add Ti in an amount of 0.005% or more. More preferably, the amount of Ti is made 0.010% or more. On the other hand, when Ti is added excessively, coarse TiC precipitates and toughness may decrease, so the addition amount is preferably made 0.15% or less. A more preferable Ti amount is 0.10% or less, and a still more preferable upper limit is 0.08% or less.

Zr: 0.005 to 0.15%
Zr has a stronger affinity for N than Ti, and in order to enhance the effect of B, it is preferable to add 0.005% or more. More preferably, 0.015% of Zr is added. On the other hand, excessive addition of Zr results in formation of coarse oxides, which may impair toughness. Therefore, the addition amount is preferably made 0.15% or less. From the viewpoint of stability of toughness, the more preferable amount of Zr is 0.10% or less, and the more preferable amount of Zr is 0.08% or less.

  Furthermore, in the present invention, in order to control the form of inclusions such as oxides and sulfides, it is preferable to add one or more of Ca, Mg, Y, Ce and La.

Ca, Mg: 0.0003 to 0.0050%
Y, Ce, La: 0.0100 to 0.0050%
Ca, Mg, Y, Ce, La are elements used to control the morphology of sulfides, and one or more of them may be added to suppress the decrease in hot workability and toughness due to MnS. preferable. In particular, in order to prevent the formation of MnS stretched in the rolling direction at the center of the plate thickness, it is preferable to add Ca and Mg at 0.0003% or more, and Y, Ce and La at 0.010% or more. On the other hand, Ca, Mg, Y, Ce and La are also strong deoxidizing elements, and when added in excess, oxide clusters may be formed to lower the toughness, so 0 for Ca and Mg, respectively. It is preferable to set the content to not more than .0050% and to not more than 0.0050% for Y, Ce, and La. More preferably, Ca, Mg is 0.0040% or less, Y, Ce, La is 0.0300% or less, and Y, Ce, La is more preferably 0.0200% or less.

  In the present embodiment and the other embodiments described above, the balance other than the above-described elements substantially consists of Fe, and it is possible to add a trace amount of elements that do not harm the effects of the present invention, including unavoidable impurities. .

Next, a hot working method in the method of manufacturing a structure of the present invention will be described.
Specifically, a steel piece having the above-described chemical composition is heated to 900 ° C. to 1100 ° C. and hot-rolled to form a steel plate having a thickness of 5 mm or more. After that, the structure of the present invention can be manufactured by applying a heat treatment before welding and welding, and a heat treatment after welding under the conditions as described above.

The present invention is characterized by containing Ta. As described above, Ta precipitates at grain boundaries to improve the grain boundary coverage, and contributes to the improvement of creep strength of the material as precipitate series after grain boundary movement in long-term creep deformation. Moreover, this property also works effectively to prevent Type IV damage, and by preventing coarsening together with the M ₂₃ C ₆ type carbide, the weld heat affected zone can be made completely equivalent to the base metal. However, Ta has the property of precipitating as a nitride at high temperatures, particularly at preheating (usually 1100 to 1300 ° C.) of hot rolling. Therefore, in a material containing a large amount of Ta, care must be taken to prevent hot working cracking (rolling cracking) due to grain boundary precipitation of TaN. This phenomenon was first found at the development stage of the present invention.
Heretofore, it has not been found at all that the addition of Ta lowers the hot workability. Therefore, in order to ensure the hot workability, the maximum heating temperature before hot rolling is lowered to 1100 ° C. or less to prevent the coarsening of TaN. On the other hand, the reduction of the heating temperature before the hot rolling raises the strength during rolling of the steel material and raises the roll reaction force at the end of the rolling. For this reason, the hot rolling itself becomes difficult, and in some cases, the situation that the hot rolling can not be performed (the mill stiffness can not follow and can not be reduced) occurs. Therefore, management of the hot rolling end temperature is also important.

FIG. 11 is a diagram (graph) for schematically explaining the hot rolling temperature and the state of the steel plate at the time of hot rolling in a steel slab having a composition within the range of the present invention, and the vertical axis is measured Shows the surface temperature of the steel sheet after each rolling pass at the time of hot rolling on the horizontal axis, and in the graph, the state of the surface of the steel sheet (with hot working cracks ●, no cracks ○, hot working Cracked (not hot-rolled) ×) was shown.
When the temperature during hot rolling is 900 ° C. to 1100 ° C., no cracking occurs on the surface of the steel plate, and if it is higher than 1100 ° C., working cracking (grain boundary precipitation embrittlement of TaN) occurs, more than 900 ° C. When it is low, it can be understood that the test can not be performed because the rolling can not be performed. That is, when manufacturing by hot rolling, it turns out that it is necessary to make extraction temperature 1100 ° C or less and end temperature of rolling 900 ° C or more. If the temperature deviates from this range, it is understood that it is necessary to cool the steel plate (stand by for rolling) or to reheat it in the heating furnace.
Note that FIG. 11 shows the results of a small scale hot rolling machine of laboratory scale, and the absolute value of the roll reaction force has no relation with the actual machine manufacturing apparatus, but the roll diameter, the plate width and the plate thickness Since the ratio to the rigidity of the mill is matched to the actual machine at the same time, the roll reaction force under rolling conditions has a function capable of reproducing the hot working crack of the actual machine as a ratio to the maximum reaction force. From this ratio, the steel plate is cracked at 80% of the maximum roll reaction force in high temperature rolling (1100 ° C or more), and the mill rigidity is not the cause, and it may exceed 100% at low temperature. It can be interpreted that the hot rolling could not be performed due to the rolling load exceeding the capacity.

  As mentioned above, although the high strength ferritic heat resistant steel structure concerning this embodiment and its manufacturing method were explained, in the present invention, the effect of the present invention can fully be exhibited, without limiting the groove shape in particular. That is, in addition to the "V groove" described in the present embodiment, the "X groove", "I groove", "K groove", etc. may be appropriately selected depending on the application, size, and the like.

Hereinafter, the present invention will be described in more detail by way of examples.
The conditions in the following examples are one condition example adopted to confirm the practicability and effects of the present invention, and the present invention is not limited to this one condition example. The present invention can adopt various conditions as long as the object of the present invention is achieved without departing from the scope of the present invention.

  Steels having the chemical components shown in Tables 3 and 4 were melted and cast into ingots by melting with an electric furnace or by a manufacturing process having a molten steel consistent process. Subsequently, this ingot is preheated to a pre-rolling pre-rolling temperature shown in Table 5 and then formed into a steel pipe or steel plate having a required shape by hot rolling or hot forging at a hot working finish temperature shown in Table 5 Hot processed. Thereafter, "normalizing" at a temperature range of 1000 ° C. to 1180 ° C. to form a martensitic structure in the steel, and subsequently tempering at a temperature range of 720 ° C. to 760 ° C. for 1 hour or more. "Tempered martensite" single phase and simultaneously soften the material, then apply heat or cold working to reheat the steel pipe or steel plate to a temperature below the tempering temperature to obtain the desired final shape Obtained. Table 5 shows the final shape (“steel pipe” or “steel plate”) of the obtained structure.

The welding groove was processed into the steel pipe or steel plate which these processing completed, and "the heat treatment before welding" was performed on the conditions of Table 5. The groove used was V-groove, and the angle of the groove surface was 15 ° on one side, and the abutment portion of the route was 1 mm in thickness.
After that, the welding heat input is about 0.5 to 5.0 kJ / mm, and GTAW (gas shield system coated arc welding using tungsten electrode) or SMAW (hand bar type coated arc welding), SAW (coated arc welding) Welded using
Subsequently, post weld heat treatment was performed on the joint under the conditions shown in Table 5. Incoloy 800 (registered trademark) Ni-based alloy was used as the weld metal.

  In order to measure the creep properties of welded joints, a flanged round bar test piece with a diameter of 6 mm in parallel and a length of 30 mm in parallel is oriented perpendicular to the welding direction so that the heat-affected zone of the welded joint is in the parallel. In a direction perpendicular to the thickness direction of the steel pipe or steel plate, and subjected to a creep test at 700 ° C. for 1 thousand hours (a creep test equivalent to 100,000 hours corresponding to 650 ° C., an assumed environment to which the welded joint of the present invention is exposed Conditions were assumed to give the material a structural change substantially equivalent to

Similarly, for a base steel, a 2 mm V-notch Charpy impact test described in JIS Z2242 was carried out at room temperature on a 10 mm square, 55 mm long test piece with a 45 ° V groove. Since the base steel requires an absorbed energy of 27 J or more at room temperature as a standard for processing as a structure, its threshold value is 27 J. The results are shown in Table 6 as BCH (J).
In addition, in order to evaluate the toughness of the weld heat affected zone interface of a welded joint, according to the method and test piece shape described in JIS Z2242 as in the case of the base steel, the test piece is sampled from the thickness center position and the absorbed energy at room temperature It was measured. The notch is processed so that it is horizontal to the plate thickness direction with the point at which the interface between the weld joint metal and the weld heat affected zone intersects the line at the plate thickness center (plate thickness center position with reference to the interface). . The threshold value was set to 20 ° C., 27 J, which is the minimum absorbed energy at which cracking does not occur during processing of the thermal power generation member. The results are shown in Table 6 as WCH (J).

  In addition, for the properties of the base steel, creep test specimens of the same shape as in the case of the welded joint are taken from the non-welded steel from the part where the weld groove is not formed so that all parallel parts have the structure of the base steel. The creep test is conducted and evaluated under the same test conditions as the welded joint. The fact that the base steel has a creep strength of 80 MPa or more at 650 ° C. for 100,000 hours can be confirmed by the temperature accelerated creep test at 700 ° C. for 1,000 hours, so the creep strength threshold value of the base steel was 80 MPa. The results are shown in Table 6 as BCR (MPa).

Furthermore, the residual γ amount of the structure was measured by the X-ray diffraction peak height calibration curve method. In addition, 10 points are measured to include the weld metal, HAZ, and base material in the direction perpendicular to the weld line from the groove, and 5 points or more, that is, 50% or more of the length of the steel pipe or the width of the steel plate The case where the amount of residual γ was 0.2% by volume or less was judged as good.
The results are shown in Table 6 as RAA. "(Circle)" in a table | surface shows that it was favorable and "x" was defect.

After heat treatment after welding, it was cut at a cross section perpendicular to the weld line and metal flow corrosion was performed to reveal HAZ. After taking a steel sample from HAZ and processing it into a sample for electron microscopy, the occupancy (coverage) of M ₂₃ C ₆ system carbides and TaC at grain boundaries was measured using a scanning electron microscope. In this case, the specification of M ₂₃ C ₆ system carbides and TaC was carried out by pattern analysis of reflection peaks by X-ray diffraction of EDX (electron beam scattering X-ray spectrometer) and electrolytic extraction residue. In addition, before observation with a scanning electron microscope, the angles of adjacent conglomerates in the grain boundary are previously analyzed by EBSD (electron beam backscattering diffractometer), and the coverage by M ₂₃ C ₆ system carbide and TaC is It was confirmed that the grain boundaries being measured were large angle grain boundaries, in particular, the angles around the 110 common rotation axis were 60 ° and 54 °. This is because the 16 ° grain boundaries could hardly be confirmed in this test.

Further, the inter-precipitate distance on the large angle grain boundary is measured on the image in the manner shown in FIG. 2 from the observation result of the M ₂₃ C ₆ system carbide and TaC obtained by the above-mentioned microscopic observation, The distance was approximated easily as the following formula (2), assuming that the particles had a square distribution on the boundary surface.

λ _ave = 1.4 [l _s ] _ave − [d _s ] _ave (2)
Here, [l _s ] _ave is the mean particle center-to-center distance (nm), and [d _s ] _ave is the mean particle diameter (nm). λ _ave is the average particle surface distance (nm).

Table 7 shows the essence of the present invention, the circle equivalent average particle diameter R (nm) of M ₂₃ C ₆ system carbide and TaC deposited on the large angle grain boundary of the welding heat affected zone, the same on the large angle grain boundary Length occupancy ratio P (%) by precipitate, distance DP between particles surface on the large angle grain boundary of the same precipitate (nm), creep strength of weld joint at 700 ° C for 1 hour and creep rupture strength of base material ( The ratio of BCR) is expressed as CRHAZ, expressed as% at 100%. Although CRHAZ theoretically does not exceed 100%, it may be possible as a range of variation of estimated values if it exceeds 100% because of the accuracy of the temperature accelerated test.

In the present invention, when it is judged that the effect of the present invention is exhibited, temperature-accelerated creep test of weld joint (temperature-promoted creep test in the name of creep test equivalent to 650 ° C. for 100,000 hours at 700 ° C. for one thousand hours) It is also considered that the result does not fall below 5% which is the allowable variation range of the creep strength of the base material specified by ASME standard, ASTM standard or the like.
In addition, although the fracture surface observation of the creep rupture test piece was actually carried out simultaneously and the occurrence of Type IV damage was also confirmed at the same time, the strength in the temperature accelerated creep test of the creep rupture strength of the welded joint which passes the above judgment criteria. When it was within 5% of the creep rupture test strength of the base material, it was confirmed that Type IV damage did not occur. That is, it was confirmed that a creep void connected low ductility fracture surface was not generated along the HAZ outer edge.
In Table 5, the conditions for heat treatment before welding applied to welded joints (width or length of one side from the center of the welded groove, heat treatment temperature, heat treatment time) and heat treatment conditions after welding (heat treatment temperature after welding, after) The heat treatment time was described. Furthermore, the hot rolling conditions (pre-rolling preheating temperature, hot working finish temperature) are also described.

Tables 8 and 9 show the chemical components of the comparative steels, and Tables 10 and 11 show the weld joint manufacturing conditions and the heat treatment conditions before welding, the heat treatment conditions after welding, the hot working conditions (preheating temperature before rolling, hot The processing end temperature) is shown as in Table 5.
Of the chemical compositions of the M ₂₃ C ₆ -based carbides observed on the large angle grain boundaries in the fine grain HAZ of the present invention steel having the chemical components shown in Tables 3 and 4, the constitution of M is substantially 70 %, At the time of preparing a creep test piece for evaluation of a welded joint separately that at least 1% consists of Cr, Fe, W, a test piece for TEM observation is also collected, and EDX analysis (magnification of 10,000 times) ( It confirmed together with the energy dispersive X-ray analyzer. This composition may deviate when the post-welding heat treatment conditions according to the present invention are not satisfied, and in particular, when M ₂₃ C _6- based carbides mainly composed of Mn are formed, their coarsening is extremely fast. Therefore, it should be noted that the creep rupture strength of the welded joint is lowered.

In Table 12, similarly to Table 6, the creep rupture strength BCR (MPa) of the base material, the Charpy absorbed energy WCH (J) before the creep test of the weld heat affected zone interface of the welded joint, the Charpy absorbed energy BCH of the base material J) An evaluation result RAA of the residual γ amount of the structure is shown. In Table 13, as in Table 7, the circle equivalent average particle diameter R (nm) of M ₂₃ C _6- based carbides precipitated on the large angle grain boundaries of the weld heat affected zone, the length by the same precipitate on the large angle grain boundaries The percentage of area occupancy (coverage) P (%), the ratio of the creep strength of a weld joint at 700 ° C. for 1 thousand hours to the creep rupture strength of a base metal in 100%, is shown as CRHAZ.

Among the comparative weld joints, 51 is an example in which the C concentration of the base steel is insufficient as 0.01% or less and the creep strength of the base steel itself is lowered, and 52 is because C contained over 0.12%. This is an example in which the creep rupture strength of the steel, including the welded joint, is lowered because the equivalent circle diameter in the electromicroscopy observation of the M ₂₃ C ₆ type carbides is coarsened to exceed 300 nm.

53 is an example in which the addition amount of Si is lower than 0.02%, deoxidation is insufficient, coarse oxide clusters are formed, and the toughness of the base steel and the weld joint is lowered, and 54 is Si This is an example in which the addition amount is excessive, the precipitation of Fe ₂ W is promoted, and the creep strength of the base steel is lowered.

  55 is an example in which the addition amount of Mn is insufficient, deoxidation is insufficient and oxide clusters are formed in the base steel, and the toughness of the base steel is lowered, and 56 is excessive addition of Mn This is an example in which the creep strength of the base steel decreases.

Since 57 and 58 contained the impurity elements P and S, respectively, beyond the scope of the present invention, they are both examples in which the creep strength of the base steel is lowered.
No. 59 is an example where the content of Ni as an impurity is excessive and the creep strength of the base steel is lowered.
No. 60 is an example in which the Cu content as an impurity is excessive, causing red shortness during hot working of the base steel, and the toughness being lowered.

  61 is an example in which the addition amount of Co is insufficient, the δ ferrite area ratio rises excessively, and the toughness of the base steel is impaired, and the creep strength decreases, and 62 is an excessive addition of Co, This is an example where the sigma phase precipitates everywhere and the toughness of the base steel and the welded joint is reduced.

No. 63 is an example in which the amount of addition of Cr is insufficient, the generation amount of M ₂₃ C ₆ type carbides is insufficient, and the creep strength of the base steel is lowered, and at the same time the coverage of precipitates on the large angle grain boundary of HAZ is lowered. In contrast, 64 is an example in which the creep strength of the base steel is lowered because the Cr addition amount is excessive to promote coarsening of the M ₂₃ C ₆ type carbide.

  In the present invention 65, Mo, which should be intentionally avoided in the present invention, is inevitably mixed as an impurity, and the content thereof exceeds the upper limit of the present invention, so the formation and growth of carbides are accelerated and coarsening progresses As a result, this is an example in which the creep rupture strength of the base steel is lowered.

66 is an example in which the content of W, which is a feature of the present invention, is insufficient and the creep strength of the base steel is insufficient, and 67 is an excessive addition of W to promote coarsening of Fe ₂ W, and the base steel Is an example in which the creep strength of

  68 is an example in which the Nb addition amount is less than the lower limit of the present invention, the precipitation strengthening by NbC does not work effectively, and the creep strength of the base steel decreases, and 69 is because the Nb amount exceeds the upper limit of the present invention Also, this is an example in which coarse precipitation of NbC occurs and the precipitation strengthening does not function in the same manner, and the creep strength of the base steel decreases.

  70 is an example where the V addition amount is insufficient and the precipitation strengthening by VN is insufficient, and the creep strength of the base steel decreases, and 71 is that the V addition amount is excessive and coarse VN precipitates, and the base steel This is an example in which the toughness is reduced.

  In the present invention, 72 is precipitated on the large angle grain boundaries and insufficient in Ta which exerts an effect on Type IV prevention and creep strength improvement of the base steel, and the creep rupture strength of the base steel decreases and the large angle grain of HAZ This is an example in which the grain boundary coverage by precipitates in the world is less than 40% and Type IV damage occurs. On the other hand, 73 is an example in which a steel plate test piece can not be produced because the amount of added Ta is excessive and the TaC at grain boundaries is coarsened, the hot workability is lowered, and the working crack is generated.

  74 is an example in which the addition of Nd is insufficient and the effect of alleviating impurity segregation is insufficient, and the toughness of HAZ is lowered. On the other hand, 75 is excessive in Nd addition to form NdS, and base steel and HAZ Is an example where the toughness of

76 and 77 because the amount of Ti or Zr exceeds the upper limit of the present invention, respectively, both precipitated coarse TiC and Zr ₂ O _3, or mother steel and HAZ toughness are examples of reduced was crystallized .
78 and 79 are examples in which the addition amounts of Ca and Mg are excessive and CaO or MgO is generated, respectively, and the toughness of the base steel and the HAZ is reduced.
In the examples 80, 81, and 82, the addition amounts of Y, Ce, and La are excessive, respectively, and all the oxide clusters are formed to reduce the base steel and the HAZ toughness.

  83 is an example in which the upper limit was exceeded and the creep strength of the base steel decreased because Al, which is an impurity of the present invention, was mixed in from the refractory during the steel making process.

  84 is an example in which the precipitation amount of nitride, VN, is reduced due to the insufficient addition amount of N, and the creep strength of the base steel is decreased. Since 85 is an excessive amount of N, BN is precipitated, This is an example where reinforcement by B, which is an important constituent element of the present invention, is not achieved and creep strength is lowered and Type IV damage occurs at the heat treatment width edge of a locally heat treated joint. This Type IV damage is also an example showing that a problem occurs when the local heat treatment is applied to a low N steel of a specified amount or less according to the present invention. Therefore, the above-mentioned large-angle grain boundary coverage P of HAZ and the creep rupture strength ratio CRHAZ of the joint / base material in Table 6 show the results of the local heat treatment width outer edge only in this example.

86 is an example in which many oxide clusters are formed and the toughness of the base steel is lowered because the impurity O exceeds the upper limit of the present invention.
No. 87 was an example where the amount of B required to improve creep strength was small and the creep strength of the base steel decreased, 88 was too much B, and weld metal itself could not be welded because weld metal high temperature cracking occurred during welding It is an example.

In Example 89, the heating temperature before rolling is less than 900 ° C., and the rolling finish temperature is inevitably less than 900 ° C., and the hot workability is lowered to cause working cracks, so that the steel plate test piece can not be prepared. 90 is an example in which the pre-rolling heating temperature was raised to over 1100 ° C., so coarsening at the grain boundaries of TaN was remarkable, and working cracks occurred during hot rolling, and a steel plate could not be produced.
In 91, although the heat treatment temperature before rolling was in the range of the present invention, since the rolling completion temperature was lower than the lower limit of 900 ° C. of the present invention, cracking occurred on the steel plate surface by rolling at 900 ° C. or less. It is an example.

No. 92 has a heat treatment temperature before welding lower than the temperature range according to the present invention, and sufficient dissolution and dissolution of carbides is not realized in the joint before welding, and the average particle diameter of M ₂₃ C ₆ type carbides in HAZ exceeds 300 nm, This is an example in which the precipitate spacing on the large angle grain boundary exceeds 150 nm, Type IV damage occurs, and the creep rupture strength of the joint is lower than the creep rupture strength of the base steel. On the other hand, 93 was an example where the heat treatment temperature before welding was too high, so all the weld heat affected zone had a two-phase structure of δ ferrite and martensite, and the creep strength of the welded joint decreased and at the same time the toughness of the welded joint also decreased. is there.

The heat treatment time before welding 94 is short, sufficient carbide decomposition and dissolution is not realized in the joint before welding, and the average grain size of M ₂₃ C ₆ type carbides in HAZ exceeds 300 nm, and precipitation on large angle grain boundaries Distance between objects exceeds 150 nm and the grain coverage by precipitates at the large angle grain boundary of HAZ is also reduced. As a result, Type IV damage occurs and the creep rupture strength of the joint is lower than the creep rupture strength of the base steel Example.

Although the temperature and time of the heat treatment before welding were appropriate, the working width was 20 mm or less, and the outer edge of the HAZ was not heated to a sufficiently high temperature (that is, less than 1000 ° C.). The remaining portion remains as it is in the joint, and coarsening occurs due to the heat history of the weld heat affected zone of the non-solid solution carbide, and the precipitate spacing on the large angle grain boundary exceeds 150 nm, Type IV damage In this example, the creep strength of the joint is reduced. In this case, the average particle size of the M ₂₃ C ₆ type carbide and TaC is greater than 300 nm, also the grain boundary coverage by the high angle grain boundary precipitates did not reach 40%.
In the example 96, the width of the heat treatment before welding is as large as 290 mm on one side and beyond the present invention, resulting in the overall heating of the target steel pipe of 500 mm in diameter. At this time, the entire structure of the member is re-quenched by pre-welding heat treatment including the base material portion, and the post-welding heat treatment is subsequently performed, but retained austenite remains partially, and the base steel This is an example in which the toughness is partially reduced. In addition, in the seam welding of a steel pipe having a seam welded portion parallel to the axial direction of the pipe, the application width of the heat treatment before welding is too wide in the seam welding of the steel pipe having a seam welded portion parallel to the axial direction.

The heat treatment temperature after welding is as low as 720 ° C. or less, the grain boundary coverage of HAZ with M ₂₃ C ₆ type carbide is less than 40%, and the precipitate spacing on the large angle grain boundary exceeds 150 nm. In the example where damage occurred and the creep strength of the joint decreased with respect to the base steel, 98 is the average diameter in the electromicroscopy observation of the M ₂₃ C ₆ type carbide of HAZ because the heat treatment temperature after welding exceeds 760 ° C. Is larger than 300 nm, the precipitate spacing on the large angle grain boundary is greater than 150 nm, and the precipitation density on the grain boundary is also reduced, and the grain boundary coverage is less than 40%. It is an example which generate | occur | produced and the creep strength of the joint fell with respect to the base steel.

  The heat treatment time after welding 99 is short, high residual stress remains in the welded joint, hydrogen invading from the weld metal is concentrated in the high residual stress area, and many cracks from the welded joint are generated the day after welding. This is an example in which a weld groove could not be formed as a result (hydrogen delayed fracture occurred).

  Although 100 is an example which applied this art to circumference welding at the time of connection of steel pipes, as a result of performing heat treatment to all the steel pipes connected before welding, a large amount of residual gamma remains in a base material, 0. This is an example in which it can not be applied as a welded structure, since the reduction of the estimated creep rupture strength of the base material and the reduction of the base material toughness are simultaneously caused because the value largely exceeds 2%.

1 Weld metal 2 coarse grained area or coarse grained area HAZ
3 fine grain area or fine grain area HAZ
4 Two-phase area or two-phase area HAZ
5 Base metal, or base metal part or base steel 7 Large angle grain boundary 12 Weld groove surface 13 Intersection (or line) of weld groove surface and base material surface
14 Minimum width for heat treatment before welding 15 Maximum width for heat treatment before welding

Claims (4)

A high-strength ferritic heat-resistant steel structure composed of a base material with a thickness of 5 mm or more, a welding heat affected zone, and a weld metal,
The composition of the base material is, by mass%, C: 0.01 to 0.12,
Si: 0.02 to 0.45,
Mn: 0.20 to 0.60,
Cr: 8.0 to 12.0,
W: 2.00 to 7.50,
Nb: 0.02 to 0.10,
Ta: 0.10 to 1.00,
V: 0.10 to 0.50,
N: 0.003 to 0.015,
B: 0.008 to 0.050,
Co: 0.50 to 3.00,
Nd: 0.005 to 0.05
Contains
Mo ≦ 0.030,
Ni ≦ 0.10,
Cu <0.05,
Al ≦ 0.005,
P ≦ 0.020,
S ≦ 0.010,
O ≦ 0.010
And the balance consists of Fe and unavoidable impurities ,
Before SL (the M and total Cr, Fe, and W, 70% or more) M ₂₃ C ₆ type carbides precipitated in large angle grain boundaries on the weld heat affected zone and the average value of equivalent circle diameter of the TaC is 300nm or less the large angle surface distance of M ₂₃ C ₆ type carbides and TaC precipitated in the grain boundary on the at 150nm or less in the Ku H AZ relation to the particle type, the large angle grain by the M ₂₃ C ₆ type carbides and the TaC A high-strength ferritic heat-resistant steel structure characterized in that the area coverage is at least 40%. Furthermore, the composition of the base material is, by mass%,
Ti: 0.005 to 0.15%,
Zr: 0.005 to 0.15%
The high-strength ferritic heat-resistant steel structure according to claim 1, wherein one or two of them are contained. Furthermore, the composition of the base material is, by mass%,
Ca: 0.0003 to 0.0050%,
Mg: 0.0003 to 0.0050%,
Y: 0.0100 to 0.0050%,
Ce: 0.0100 to 0.0050%,
La: 0.0100 to 0.0050%
The high-strength ferritic heat-resistant steel structure according to claim 1 or 2, wherein one or more of them are contained. It is a manufacturing method of the high strength ferritic heat resistant steel structure according to any one of claims 1 to 3,
The steel piece of the composition as described in any one of Claims 1-3 is heated to 900 degreeC-1100 degreeC, and it hot-rolls, and sets it as a steel plate of 5 mm or more in plate thickness,
Then, the steel sheet is subjected to groove processing, and a portion of 20 mm or more and 100 mm or less in the direction orthogonal to the weld line from the position where the groove surface contacts the base material surface, and the width or length of the heat resistant steel structure The portion which is 50% or less of the above is locally heated and maintained at a temperature of 1000 ° C. or more and 1100 ° C. or less so as to be a heat treatment time Ts or more for 30 mm or more determined by the following equation (1) And let it cool down,
Next, after welding the groove, the region between 20 mm and 100 mm is heated to a temperature zone of 720 ° C. or more and 760 ° C. or less in the direction perpendicular to the weld line from the position where the groove face contacts the base material surface. A method for producing a high-strength ferritic heat-resistant steel structure, characterized in that the temperature is maintained for 1 hour or more and then allowed to cool.
Ts = 30 × (t / 25) ^1.2 +26 (1)
However, t is a plate thickness (mm) of a base material, and Ts is a total (minute) of heating time at a heating set temperature (1000 ° C. to 1100 ° C.) of a groove before welding. In addition, t shall be 5 mm or more.

Images