JP6471287B2 - W-Cr-based alloy or mold, electrode or extrusion die produced thereby - Google Patents

Description

  The present invention relates to a die-casting die used for manufacturing various parts made of a W—Cr base alloy or an aluminum alloy using the same, a hot extrusion die such as brass, a peripheral member of a lens die, S45C and stainless steel. The present invention relates to a tool such as a base electrode used in electric heating forging.

  Die-casting dies used for manufacturing various parts made of aluminum alloys, hot extrusion dies for rods such as brass, peripheral members of dies for glass lens molding, pedestal used for electric heating forging for S45C, stainless steel, etc. Materials used for electrodes and the like are required to have excellent mechanical properties and oxidation resistance at high temperatures, and new W-based alloys invented by the present inventors are used (Patent Documents 1 to 3). ).

JP 2002-275570 A JP 2007-270339 A JP 2011-246797 A

Max Hansen, Kurt Anderko: Constitution of Binary Alloys, Second Edition, McGraw-Hill Book Company, Inc. , 1958, p. 570-571 American Ceramic Society: Phase Equilibria Diagrams Database, Version 4.0, 2014, 6462-A-EC-162-A American Ceramic Society: Phase Equilibria Diagrams Database, Version 4.0, 2014, 4229-A American Ceramic Society: Phase Equilibria Diagrams Database, Version 4.0, 2014, 10921-B American Ceramic Society: Phase Equilibria Diagrams Database, Version 4.0, 2014, 5067-A, 11671-A Max Hansen, Kurt Anderko: Constitution of Binary Alloys, Second Edition, McGraw-Hill Book Company, Inc. , 1958, p. 684-691 Max Hansen, Kurt Anderko: Constitution of Binary Alloys, Second Edition, McGraw-Hill Book Company, Inc. , 1958, p. 1024-1026 Japan Iron and Steel Institute Edition: Third Edition Steel Handbook Volume I Fundamentals, Maruzen Co., 1981, p. 6-7

  The peripheral members of the die-casting mold made of the W-based alloy of the present inventors and used for the production of various parts of aluminum alloys in Patent Documents 1 and 2 have a coefficient of thermal expansion. In addition to being small, it has excellent oxidation resistance, so there are few occurrences of thermal cracks and it has a long service life and is popular. However, the W-based alloys of Patent Documents 1 and 2 are produced by vacuum sintering, but this requires a diffusion pump and a metal heater, which increases the cost.

The base electrode for electric heating forging made of the new W-based alloy of the present inventors used in the manufacture of various steel parts such as S45C and stainless steel in Patent Document 3 is 900 in addition to excellent oxidation resistance. High-temperature hardness above ℃ is higher than SKD61 of hot die steel, as well as WC-Cr ₃ C ₂ -Co-Ni cemented carbide often used hot, so it has a long service life and is popular. is there.

  Therefore, when this is used in a hot extrusion die for forming a bar material such as brass, it is difficult to be deformed at a high temperature, so that it is possible to perform extrusion with high accuracy and is popular. However, in this case, since the life form is cracked, if there is a material that is harder to break and has a high hardness and high strength, the life can be further extended, so such a material is required.

  The present invention has been made to solve the problems as described in paragraphs 0005 and 0007, and is intended to provide a new material that can be manufactured at a lower cost and is less likely to break than the above W-based alloys. It is what.

  In addition to W, the W-based alloys of Patent Documents 1 to 3 by the present inventors are made of Ni, Fe, Co, and Cr. The main component W is 75 to 98 mass%, the total amount of the three Fe groups of Ni, Fe and / or Co is 1 to 15 mass%, the Cr amount is 1 to 20 mass%, and the composition is inevitable impurities.

  Here, the total amount of Fe and Co is 0 to 30 mass% of the total amount of three types of Ni, Fe and Co. Moreover, the thing by which 10 mass% or less of W was substituted by Ta, Ti, and / or Mo is also contained. These are produced by vacuum sintering.

  First, as a comparison with the W-based alloys of Patent Documents 1 to 3, the influence of various element additions on the alloy structure was examined. In Table 1, the total amount of Ni, Fe, and Cr is larger than that in the W-based alloys (1 to 15 mass%) of Patent Documents 1 to 3 (16.2 to 17.5 mass%), and described in paragraph 0014 below. Therefore, the compound composition of the model sample in which Nb is newly added instead of Ta and the addition amount is changed is shown.

Each raw material powder is blended into the composition shown in Table 1, mixed and pulverized, and cold formed, and at low cost unlike Patent Documents 1 to 3, using an ordinary vacuum pump and carbon heater furnace, sintering at 80 kPa Ar atmosphere. Was performed at 1490 ° C. for 2 hours to prepare a 4 × 8 × 25 mm ³ test piece. The four types of model alloys produced are referred to as samples M1 to M4 below. The sintering temperature and time were set to 1490 ° C.−2 h, which is a condition that can be applied to a thick product.

  Here, Nb was selected in place of Ta because Ta shown in Patent Documents 1 to 3 shows that in the energy dispersive X-ray analysis attached to the scanning electron microscope, the Ta Lα peak and the W Lα peak are almost equal. Since they overlap, the distance dependency of the sample composition from the sample surface and the composition analysis of each constituent phase are difficult. On the other hand, Nb does not overlap, so that the composition analysis is easy.

  Furthermore, Nb and Ta are very close in chemical properties because they both belong to Group 5 in the periodic table. Also, Nb powder is about half as cheap as Ta powder and presents in the crust. Since the number of atoms is about 20 times as large, the resource problem is small.

  Table 1 shows mass%, but there may be a case where the volume ratio and atomic ratio are compared. Tables 2 and 3 show the cases of% display. At. % Is often expressed as mol%, but at. %.

  Four types of model samples were observed by SEM, and FIG. 1 was obtained. As apparent from the enlarged SEM structure in the lower part of FIG. 1, this sample has a white gray spherical particle phase (hereinafter referred to as a phase), a gray binder phase (hereinafter referred to as b phase), and a gray-black binder phase ( (Hereinafter referred to as c phase) and black particle phase (hereinafter referred to as d phase) are observed. In the model sample, b phase was observed after a phase, and c phase increased with addition of 3.9 vol% Nb, but then decreased. The d phase was slight in all.

  In FIG. 1, a subscript p is added to each of the characters a, b, c, and d. This represents an EDX analysis point of each phase. The four types of model samples, a to d phases, were quantitatively analyzed by EDX, and the results are shown in FIG.

  From the analysis result of the a phase in FIG. 2, the a phase is W and Cr is 15 at. It was found that it was a W—Cr phase that was dissolved in about%. In addition, as shown in FIG. 3 to be described later, in X-ray diffraction, W-Cr solid solution has a peak at the position of W because W is the main component, and the a phase is W-Cr solid solution. It can be understood only after analysis.

  The fact that W powder and Cr powder are mixed and sintered to produce a W-Cr solid solution phase is consistent with the W-Cr binary phase diagram of Non-Patent Document 1. From this, it was found that the sintered alloy in the present application is a W-Cr based alloy. The fact that the W phase has become a W—Cr solid solution phase is considered to be one of the reasons why the oxidation resistance of the present alloy is excellent (the oxidation resistance is shown in the oxidation increase in Table 8 described later). These are the first findings.

From the analysis result of the b phase in FIG. 2, it can be seen that the b phase is mainly composed of W and Cr and contains Ni, O, and Fe in the M1 sample. M2~M4 sample Namely, the samples were added Nb, but not observed Nb in a phase, the b-phase become Nb is observed, Nb amount of b phase as the amount added increases increases doing.

The amount of Cr solid solution in the b phase decreases with the addition of Nb, but hardly changes even when the amount added is increased. The amount of W solid solution in the b phase is 3.1 at. % Increases, but decreases as the amount added increases thereafter. Compared to these, the amounts of Ni, Fe and O are less changed. Here, it is worth noting that there is O in the b phase.

From the analysis result of the c phase in FIG. 2, it was found that the c phase is mainly composed of Ni and Cr and contains O, W, and Fe in the M1 sample. In this binder phase, Cr decreases with the addition of Nb, but hardly changes even when the amount added is increased. The amount of Ni was 6.3 at. Slightly increased to % addition, but then decreased slightly. Compared to these, the amounts of W 2 , Fe, and O change little. It is worth noting that there is O in this c phase.

From the analysis result of the d phase in FIG. 2, it was found that the d phase contains a large amount of O and is mainly an oxide of Cr or Cr and Nb. This is a result of examining the pseudo-binary system phase diagram of _Nb 2 _O 5 _-Cr 2 _{O 3} in Non-Patent Document 2, the _Nb 2 _{O 5} similar structure (Nb, _Cr) becomes _{2 O 5-X} I thought it was.

O in the b and c phases and O in the d phase, ie, the oxides of Cr and Nb, are contained as impurities contained in O contained in the raw materials used and Ar gas (purity: 99.99% or more) in the sintering atmosphere. O ₂ appears to be the source.

  O is present in the b, c and d phases, and O is not observed in the a phase. This is because the a phase is a solid phase throughout the sintering and the other three phases are not. It is suggested from the form. That is, the b and c phases are liquid phases at the sintering temperature due to their form, and become a solid phase upon cooling, and the d phase precipitates and grows as solid phase particles when the liquid phase components Cr and Nb react with oxygen. I suggest that.

Next, in order to confirm whether O in the b phase and c phase in FIG. 2 is dissolved in the binder phase or exists as an intermetallic compound or oxide, X-ray diffraction is performed, and the result is shown in FIG. Show. From FIG. 3, a peak was recognized at the position of W (◯ peak), Ni (Ｎｉ peak), hexagonal crystal or cubic NbCr ₂ (● peak).

  Here, from the structure of FIG. 2, the peak at the W position is determined to be a W—Cr solid solution (◯), and the peak at the Ni position is determined to be a Ni—Fe-based solid solution (◇). . Since some low peaks could not be identified, they were represented by X phase. The amount of X phase seems to be small and does not affect the properties.

  The X-ray diffraction intensity of the W-Cr solid solution (◯) decreases from M1 to M4 (Nb addition amount increases), but this is that W is replaced by Nb in the composition and added. In addition to this, it is considered that the amount of W solid solution in the gray binder phase is also increased by the addition of Nb.

Hexagonal or cubic NbCr ₂ (●) appears with Nb addition, that is, M2, and fluctuations due to peak positions are observed from M3 to M4, and do not necessarily increase. NbCr ₂ is considered to be contained in the b to c phases of the binder phase.

  The Ni—Fe-based solid solution (◇) does not change with the addition of Nb, which is thought to be due to the fact that the peak height also depends on the formulation composition.

Of the above phases, the NbCr ₂ phase was not observed by SEM. Therefore, it was decided to observe from FIG. 1, and the image was enlarged up to 100,000 times to obtain FIG. In FIG. 4, since no structure was observed without etching, the test piece when sintered at 80 kPa Ar atmosphere was subjected to thermal etching at 600 ° C. for 5 minutes in a reduced pressure atmosphere of 6 Pa. It is the result of having observed b phase and c phase by SEM by magnification.

From this, a small acicular white dispersed phase having a size of about 200 nm was observed in the b phase and the c phase. In FIG. 4, the dispersed phase particles that appeared in large numbers by thermal etching were determined to be NbCr ₂ phases observed by X-ray diffraction. These are the second findings. The parent phase surrounding this dispersed phase was considered to be a phase in which W, Cr, Nb , and O were dissolved in the Ni — Fe-based solid solution phase by EDX.

  From the above, oxides could not be detected by X-ray diffraction, but it could not be determined whether O was dissolved in the binder phase. This was thought to be due to the fact that a very small amount of crystal phase was difficult to detect by X-ray diffraction and that solid solution elements could not be detected directly. Therefore, next, a related phase diagram was examined to see if O could be dissolved in the binder phase.

  FIG. 5 is a Nb—O—W ternary phase diagram according to Non-Patent Document 3. In this ternary system, it is shown that O dissolves in the Nb—W solid solution as in the Nb—O binary system. However, solid solution of O is not shown in the ternary phase diagram of Ni—O—W (Non-Patent Document 4) and Fe—O—W (Non-Patent Document 5).

  On the other hand, according to the binary phase diagram of Fe-O (Non-Patent Document 6) and Ni-O (Non-Patent Document 7), O is 0.22 mass into the liquid phase for both Fe and Ni. % And 0.4 mass%, both of which are 0.1 mass% or less in the solid phase, but are shown to be dissolved or dissolved in a small amount.

Here, according to the Ellingham diagram of Non-Patent Document 8, Nb ₂ O ₅ and Cr ₂ O ₃ are not reduced by H ₂ even if they exist at 1490 ° C., but are re-sintered at 1490 ° C. It has been found that if O decreases, the O can be determined to be O dissolved in the binder phase.

Therefore, a sample of M3 composition that has been sintered at 80 kPaAr atmosphere is subjected to heat treatment at 1490 ° C.-1 h in a 100 kPaH ₂ atmosphere, and the oxygen content of the alloy before and after the heat treatment is determined by LECO's oxygen-nitrogen simultaneous analyzer TC-500 (hereinafter referred to as “oxygen-nitrogen analyzer”) It was analyzed with an oxygen-nitrogen analyzer.

As a result, O was 0.30 mass% before the heat treatment, but decreased to 0.060 mass% after the heat treatment. As described above, since Nb ₂ O ₅ and Cr ₂ O ₃ are not reduced by H ₂ , it was determined that this reduced O was O dissolved in the binder phase.

  From the above, O recognized in the b phase is composed of O dissolved in the binder phase composed of Fe, Ni, W, and Cr in the M1 composition, and Fe, Ni, W, Cr, and Nb in the M2 to M4 compositions. It was determined that the O was dissolved in the bonded phase. O observed in the c phase is also considered to be O dissolved in the binder phase. These are the third findings.

In summary, the a phase observed in the model alloy is W-Cr solid solution particles, the b phase is Fe-Ni-W-Cr solid solution bonded phase in which O is solid solution without addition of Nb, and O is solid solution with addition of Nb. Fe-Ni-W-Cr-Nb solid solution bonded phase in which NbCr ₂ is dispersed, c phase is Fe-Ni-W-Cr solid solution bonded phase in which O is solid-solved when Nb is not added, O is solid-solved when Nb is added, and NbCr ₂ dispersed Fe-Ni-W-Cr- Nb solid solution binding phase, d-phase is at _Cr 2 _O 3, Nb added in Nb not added was regarded as _{(Nb, Cr) 2 O 5} -X.

  The types of components of the b phase and the c phase are the same, but the amounts are different, the b phase is rich in Cr, W, and Ni, and the c phase is Cr and Ni rich.

  Of the above compositions, the effect of O was not well understood. Furthermore, since Ar atmosphere sintering at 80 kPa produces d-phase, that is, oxides of Cr, Cr, and Nb, the possibility of being unsuitable as a sintering method for this alloy was considered.

In order to confirm these, a sample similar to the others was prepared as hydrogen atmosphere sintering (hereinafter referred to as 100 kPaH ₂ atmosphere sintering) in which the sintering atmosphere was heated and reduced in a hydrogen stream under atmospheric pressure. SEM observation was performed about the produced sample, and FIG. 6 was obtained. The composition is the same as each of the samples M1 to M4 in Table 1, but the sintering method is different, so the sample symbols are shown as M1 ′ to M4 ′.

In the sample sintered at 100 kPaH ₂ atmosphere, phases corresponding to the a phase and the b phase in FIG. 1 were observed in all the samples, and are shown in FIG. 6 as a ′ phase and b ′ phase, respectively. Further, the phase corresponding to the c phase in FIG. 1 is shown in FIG. 6 as the c ′ phase. In addition, although c 'phase was recognized by the 0-7.9 vol% Nb addition sample, it was not recognized by the 11.9 vol% Nb addition sample.

A phase corresponding to the d phase in FIG. 1 was hardly observed in any of the samples. This is because Ar containing impurities as O ₂ was not used, and oxide contained in the raw material was reduced and removed during sintering by 100 kPaH ₂ atmosphere sintering, and as a result, almost no oxide was generated in the sintered sample. Indicates no.

  FIG. 7 shows the results of analyzing the a ′ phase, b ′ phase, and c ′ phase of FIG. 6 by EDX attached to the SEM. The a 'phase had substantially the same W and Cr amounts as the a phase. In the b 'phase, the amount of O decreased compared with the b phase and became slight. Therefore, the amount of Cr and Ni was slightly larger than the b phase. O is not recognized in the c ′ phase, and it is thought that Ni is more than the c phase.

  Compared to the samples M1 to M4 in FIG. 1, the samples M1 'to M4' are atomized. Initially, it was estimated that the amount of W solid solution in the binder phase increased. However, as shown in Table 4, the results of examining the area ratio of phases other than the W-Cr phase by image analysis are shown in Table 3 It was found that the area ratio values in M4 and samples M3 ′ to M4 ′ were almost the same in the corresponding samples.

  Therefore, the expectation of “because the amount of W solid solution in the binder phase increased” was completely denied. From this, the reason why the samples M1 ′ to M4 ′ are atomized is that the solid solution amount of O in the binder phase is decreased, and the solid solution amount of Cr is increased correspondingly. It was thought that the grain growth of the -Cr phase (particles) was suppressed.

  As a result of analyzing the oxygen content of the alloy of M4 and M4 'with an oxygen-nitrogen analyzer, it was confirmed that the former was 0.26 mass%, whereas the latter was 0.026 mass%. The amount of O by MDX b phase EDX was about 3 at. %, Which corresponds to 0.6 mass%, the amount of the binder phase is about 50 vol%, and considering the accuracy of EDX, it can be said that the results are almost the same as the results obtained by the oxygen-nitrogen analyzer.

  Further, the hardness of the binder phase of the b phase of M4 and the b 'phase of M4' was measured by micro Vickers and compared. This is because in the measurement of the hardness of the entire alloy, the size of the a-phase and a′-phase, that is, the W—Cr phase particles influences the hardness, so that the solid solution hardening amount due to O of the binder phase was not known.

As a result, the b phase showed 1680 HV (0.098 N), the b ′ phase was 1310 HV (0.098 N) , and the former was confirmed to be solid solution hardened (oxygen solid solution strengthening) by O. This is the fourth finding. In the case of Nb added is dispersion-hardened by NbCr _2.

FIG. 8 shows the measurement result of M3 ′ by X-ray diffraction together with the result of M3 described above. The amount of NbCr ₂ was not significantly different between M3 ′ and M3. The reason why (Nb, Cr) ₂ O _5-X is not recognized even in M3 ′ is because the amount is small as described in paragraph 0036. In addition, it can be seen that (Nb, Cr) ₂ O _5-X is small in amount when there is no peak and when combined with the oxygen analysis result in paragraph 0053.

  Note that the sample M4 'was finely crushed during the grinding process, so that it was found that this alloy was brittle when the amount of O dissolved was small. Here, the cause of embrittlement is considered. First, it can be considered that embrittlement is caused by an increase in the solid solution amount of Cr instead of the solid solution of O in the binder phase.

  However, in this case, the amount of solid solution of Cr is smaller than that in the case where Nb is not added, so that it seems that there is little embrittlement accompanying the increase of the amount of solid solution of Cr. Therefore, it was considered that the oxygen solid solution strengthening clarified in Paragraph 0055 disappeared, and thus embrittlement occurred. In order to obtain practical strength, O must be dissolved in the binder phase.

  The bending strength often used as a measure of embrittlement, that is, the strength, is not limited by the amount of O because it is affected by the size and distribution of the W-Cr phase, but the amount of O of the invention alloy having a useful bending strength. Is 0.051 mass% or more and 0.085 mass% or less (showing bending strength in Table 8 to be described later, and showing O amount in the alloy in Table 7), it was considered that an O amount in this range was necessary. .

The solid solution amount of O is determined by the amount of the binder phase and the atmosphere at the time of sintering, but in an Ar atmosphere of 80 kPa, the solid solution amount is insufficient when O ₂ is less than 400 ppm. If the O ₂ content is more than 600 ppm, the sinterability will be reduced and densification will not occur.

As transition metals to be added, Ta and Ti can be considered in addition to Nb. Similar characteristics were obtained even when Nb was Ta and / or Ti. From the above, Cr amount, and defining the addition amount of Nb, by performing suitable oxygen solid solution strengthening, it was found that it is possible to obtain a breakdown hardly alloys. This is the fifth finding.

As for Ta, in the same manner as Nb, oxygen solid solution strengthening was observed, and dispersion strengthening by forming TaCr ₂ as an intermetallic compound was observed. For Ti, although oxygen solid solution strengthening was observed, no intermetallic compound was formed, and no dispersion strengthening due to the intermetallic compound was observed.

Instead, Ti oxide was finely dispersed, and dispersion strengthening was observed. The reason why the oxide of Ti is finely dispersed is that TiH ₂ used as the raw material powder is more easily pulverized than Ta and Nb. The particle sizes of W, Cr, Nb, Ta, and TiH ₂ of the raw material powder to be used may be any as long as a predetermined hardness and bending strength can be obtained. This is the sixth finding.

  The above results were the same even when the 80 kPa Ar atmosphere sintering was changed to 60 or 100 kPa Ar atmosphere sintering. Here, if the pressure is lower than 60 kPa, deterioration of the furnace over time due to volatilization of Ni and Fe occurs, resulting in high maintenance costs and high costs. Further, when the pressure is higher than 100 kPa, the convective heat loss during heating becomes large, and more power is used and the cost is increased.

  It is clear that the oxidation resistance of the alloy of the present invention, in which W is W—Cr and Cr is dissolved in the binder phase, is excellent because the oxidation increase of the invention alloy and the reference alloy in Table 8 described later is smaller than that of the existing alloy. is there. Thus, the present invention has been completed.

  The composition is summarized as follows. In the present W-Cr-based alloy, the total amount of Ni and Fe is preferably 1.5 mass% or more and 3.1 mass% or less of the whole alloy. When the content is less than 1.5 mass%, the sinterability is lowered to obtain a dense alloy. It becomes difficult. When it exceeds 3.1 mass%, the high temperature hardness decreases.

  The amount of oxygen in the alloy is preferably 0.051 mass% or more and 0.085 mass% or less. When it is less than 0.051 mass%, it is difficult to obtain oxygen solid solution strengthening. If it is more than 0.085 mass%, it will be difficult to sinter and the bending strength will be lowered, making it impossible to obtain the required strength.

  In the case of peripheral members of aluminum alloy die-casting dies and glass lens molding dies, the total amount of Cr, Nb, Ta and Ti in this W-Cr-based alloy is 2.0 mass% or more and 7.6 mass of the whole alloy. % Of which Cr is preferably 2.0 mass% or more and 5.0 mass% or less.

  If the total amount of Cr, Nb, Ta and Ti is less than 2.0 mass% of the entire alloy, the hardness will be lower than 33.8 HRC and the practical hardness will be insufficient. If it exceeds 7.6 mass%, it becomes difficult to sinter and the bending strength becomes lower than 830 MPa, so that practical strength cannot be obtained. When Cr is less than 2.0 mass%, oxidation resistance is insufficient and solid solution strengthening by Cr is insufficient. If it exceeds 5.0 mass%, it becomes difficult to sinter and the bending strength becomes lower than 830 MPa, so that practical strength cannot be obtained.

In the case of a base electrode for electric heating forging of steel materials such as S45C and stainless steel, the total amount of Cr, Nb, Ta, Ti in this W-Cr based alloy is 3.0 mass% or more and 7.6 mass% or less of the entire alloy. Of these, Cr is preferably 2.3 mass% or more and 5.0 mass% or less.

If the total amount of Cr, Nb, Ta and Ti is less than 3.0 mass% of the entire alloy, the hardness will be lower than 41.8 HRC and the practical hardness will be insufficient. If it exceeds 7.6 mass%, it becomes difficult to sinter and the bending strength becomes lower than 830 MPa, so that practical strength cannot be obtained. Since Cr is used at a relatively high temperature, if it is less than 2.3 mass%, oxidation resistance is insufficient and solid solution strengthening by Cr is insufficient . If it exceeds 5.0 mass%, it becomes difficult to sinter and the bending strength becomes lower than 830 MPa, so that practical strength cannot be obtained.

  In the case of a hot extrusion die for forming a bar material such as brass, this W-Cr-based alloy has a total amount of Cr, Nb, Ta, and Ti of 4.0 mass% or more and 7.6 mass% or less of the entire alloy. Of these, Cr is preferably 2.3 mass% or more and 5.0 mass% or less.

Cr, Nb, Ta, a total amount of Ti is dispersion strengthened oxide containing intermetallic compound and / or Ti is less than 4.0 mass% of the total alloy is insufficient, the actual use strength can not be obtained. If it exceeds 7.6 mass%, it becomes difficult to sinter and the bending strength becomes lower than 1572 MPa, so that practical strength cannot be obtained. When Cr is less than 2.3 mass%, solid solution strengthening by Cr is insufficient, the actual use intensity can not be obtained. If it exceeds 5.0 mass%, it becomes difficult to sinter and the bending strength becomes lower than 1572 MPa, so that practical strength cannot be obtained.

The invention alloy according to the present application is a W-Cr-based alloy, unlike the W-based alloy of the existing alloy, and the binder phase is Cr alone or a composite with one or more of Nb, Ta, and Ti. By adding an appropriate amount, Cr is solid-solution strengthened, solid solution hardening (strengthening) with O, and / or dispersion-hardened with one or more of three kinds of oxides of NbCr ₂ , TaCr ₂ and Ti. Therefore, it is harder and stronger than conventional W-based alloys. From the above, the industrial utility value of the invention alloy is high.

It is the result of having observed SEM about the model sample which sintered the sample of M1-M4 composition at 80 kPaAr atmosphere. a _p , b _p , c _p , and d _p are positions where the respective phases, that is, the a, b, c, and d phases are spot-analyzed by EDX. It is the result of having carried out the quantitative analysis of the component element by EDX attached to SEM about the ad phase shown in FIG. ○ is W, ◇ is Ni, ◆ is Fe, □ is Cr, ● is Nb, and Δ is O. It is the result of carrying out X-ray diffraction about the test piece at the time of carrying out 80kPaAr atmosphere sintering about the sample of M1-M4 composition. The results of identifying each peak with ICDD data are shown on the figure. The X-ray type is CuKα. Since no structure was observed without etching, the test piece when sintered at 80 kPa Ar atmosphere was subjected to 600 ° C. for 5 min in a 6 Pa vacuum atmosphere, and the sample of M3 composition was subjected to b phase and c It is the result of having expanded and observed the phase by SEM. FIG. 3 is a Nb—O—W ternary phase diagram according to Non-Patent Document 3. The text size has been partially enlarged to make it easier to read. SEM observation of the model sample was performed on the test specimen obtained by heating and reducing the sample of M1 to M4 composition in the hydrogen stream under atmospheric pressure instead of the 80 kPaAr atmosphere in FIG. 1 (100 kPaH ₂ atmosphere sintering). It is the result. a ′ _{p to} c ′ _p indicate positions where each phase is spot-analyzed by EDX. The symbol corresponds to the same name in FIG. 1, but a prime (') is added to indicate that the sintering atmosphere is different. It is the result of having quantitatively analyzed by EDX attached to SEM about the a'-c 'phase shown in FIG. ○ is W, ◇ is Ni, ◆ is Fe, □ is Cr, ● is Nb, and Δ is O. And 80kPaAr atmosphere Sintered M3 in FIG. 3, for both test pieces were 100KPaH ₂ atmosphere sintering M3 ', the result of X-ray diffraction. The results of identifying each peak with ICDD data are shown on the figure. The X-ray type is CuKα. It is a figure which shows the influence of the ratio of AM with respect to BM which acts on a bending strength. The ratio of AM to BM was determined as AM ÷ BM × 100. It is a figure which shows the influence of the ratio of AM with respect to BM which acts on hardness.

  The material of the present invention can be produced by a usual powder metallurgy method. In other words, W, Ni, Fe, Cr and, if necessary, one or more of Nb, Ta, and Ti are blended into a predetermined composition, dried by wet mixing using a ball mill or attritor, and then formed into a desired shape. Compression molding is performed at a press pressure of 100 to 500 MPa.

Next, the compact is sintered at 1350-1500 ° C. for 30-120 min in an Ar atmosphere . Here it was introduced at 100kPa pressure below than 60kPa at 1350-1500 ° C. 99.9% pure Ar gas, in the form contained in the Ar gas is introduced _{O 2} as follows 600ppm or 400ppm in Ar. After sintering, the final shape is formed by cutting, grinding and / or electric discharge machining, finished, and commercialized.

A sample was prepared with the composition shown in Table 5, and each characteristic was measured. The raw material powder is a W powder having an average particle size of about 6 μm (oxygen amount 0.027 mass%), a Ni powder having an average particle size of about 2.5 μm (oxygen amount 0.15 mass%), and an Fe powder having an average particle size of about 5 μm (oxygen amount 0.02 mass%). 63 mass%), Cr powder having an average particle size of about 40 μm (oxygen amount 0.64 mass%), Ta powder having an average particle size of about 6.9 μm (oxygen amount 0.13 mass%), TiH ₂ powder having an average particle size of about 6.9 μm (oxygen) Nb powder (oxygen amount 0.23 mass%) having an average particle size of about 8.25 μm was used.

The reason for using TiH ₂ powder instead of Ti powder alone is that Ti powder easily reacts with oxygen, and low-oxygen powder cannot be obtained at low cost. TiH ₂ powder is relatively inexpensive and has high purity. This is because H ₂ dissociates into Ti during sintering.

The blended powder was mixed for 24 hours in a wet ball mill, subjected to cold forming at a press pressure of 100 MPa to 500 MPa, and sintered in a vacuum (0.8 Pa, sample symbol T1, hereinafter referred to as T1), or a purity of 99. A 9% O ₂ amount of 400 ppm or more and 600 ppm or less of Ar gas was introduced at a pressure of 60 kPa or more and 100 kPa or less, and sintered at 1490 ° C. for 2 hours (samples other than T1) 4 × 8 × 25 mm ³ test piece Obtained.

  Table 6 shows the composition expressed in vol% and shows the value obtained by dividing AM by BM and multiplying it by 100 in order to accurately grasp the effects of various additives (hereinafter referred to as the ratio of AM to BM). Here, BM indicates binder metal, specifically, the total amount (vol%) of Ni and Fe. AM refers to additive metal, and is specifically the total amount (vol%) of Cr, Ta, Ti and Nb.

Table 7 shows the types of main phases, main phases, binder phases, and oxide phases after Sintering at 1490 ° C. for 2 hours for the samples having the compositions shown in Table 6 and SEM. Quantified by the number of pixels from the image and expressed in vol%. The results of measuring the amount of O in the alloy are also shown. As already described, only the sample symbol T1 is a conventional vacuum sintering material. The other is an alloy obtained by sintering Ar gas at a pressure of 60 kPa to 100 kPa. The notation of A and AB in the invention alloy will be described in paragraph 0088.

Table 8 shows the specific gravity, bending strength, hardness, elongation, and oxidation increase (weight difference before and after holding in the atmosphere at 800 ° C. for 30 minutes), die casting mold of aluminum alloy, brass, etc. The results are shown together with the results of practical tests using hot extrusion dies. In addition, the ratio of AM to BM shown in Table 6 is shown again for the purpose of easy understanding.

Here, as shown in paragraphs 0096 to 0100, the invention alloy that can be used for the die-casting die for aluminum alloy is an A-invention alloy, which is used for electric heating forging of various steel parts such as die-casting die, S45C and stainless steel. The invention alloy that can be used for all of the base electrode and the hot extrusion dies used for forming rods such as brass was used as the AB invention alloy.

  Specific gravity was determined by the composition. The ratio of the bending strength shown in Table 8 and the ratio of AM to BM is shown in FIG. 9 as the effect of the ratio of AM to BM on the bending strength, but the bending strength shows a strong correlation with the ratio of AM to BM. I couldn't. The origin of destruction was not clear because the fracture surface was flat.

  The hardness shown in Table 8 and the ratio of AM to BM are shown in FIG. 10 as the effect of the ratio of AM to BM on the hardness. There is also a relationship between the ratio of AM to BM and the particle size of the W-Cr phase. Only a weak positive correlation.

  While the elongation of the general W-based alloy shown in Table 8 is 6%, the elongation of the sample is as small as 0.4%, which is due to solid solution hardening with O, Cr or the like. The elongation was determined by a room temperature tensile test.

The sample added with AM is excellent in that the oxidation increase at 800 ° C. for 30 minutes is 22 g / m ² or less. Therefore, it can be seen that even if an oxide layer is formed on the tool surface in practical use, it does not easily become deep. In view of cost effectiveness, the increase in oxidation needs to be 22 g / m ² or less.

  In FIG. 9 and FIG. 10, the results of M1 to M4 are omitted because the inside of the test piece was insufficiently densified and an accurate value was not obtained.

  Although the mechanical properties do not necessarily match those of the samples having similar compositions in Patent Documents 1 to 3, this is because the sintering temperature and time, the W particle size of the raw material used, and the like are different. Therefore, a strict comparison was not possible. However, as described in paragraphs 0005 and 0013, the alloy according to the present invention has an advantage of being sintered at a lower cost.

  In addition, the tendency of the high temperature hardness which was not shown in Table 8 was as follows. The invention A alloy S1 has a room temperature hardness of 390 HV (9.8 N), at a high temperature of 900 ° C., 180 HV (9.8 N), the AB invention alloy S2 has a room temperature hardness of 590 HV (9.8 N), At a high temperature of 900 ° C., it showed a hardness of 300 HV (9.8 N) and was excellent. The existing alloy T2 had a hardness at room temperature of 320 HV (9.8 N), and at a high temperature of 900 ° C., it became 140 HV (9.8 N).

  The existing alloy T1 seems to be at a practical level at first glance from the viewpoint of hardness and bending strength, but the binder phase is soft and stretched because it is not subjected to dispersion hardening (strengthening) or oxygen solid solution strengthening by intermetallic compounds. 0%, deformed with little stress when used for die casting molds for aluminum alloys, base electrodes used in electric heating forging of various steel parts such as S45C and stainless steel, or hot extrusion dies used for forming rods such as brass So it was not practical at all.

Table 8 shows the performance when samples other than T1 were used in an aluminum alloy die-casting die. T2 was unpopular because of its rough skin. B5 and C5 were cracked due to insufficient strength. △ in the die cast mold row shown in Table 8 is less rough than T2, ○ is less rough than T2, life is more than twice, and ◎ is less rough than T2. In addition, the service life is more than four times.

The results of applying samples other than T1 to hot extrusion dies used for molding rods such as brass are shown in Table 8, but the x in the row of hot extrusion dies is insufficient in strength and hardness, resulting in cost effectiveness. It was not practical. ○ indicates that the life until breaking from T2 is 3 times or more, and ◎ indicates that the life until cracking is 6 times or more.

  In the row of die-cast molds in Table 8, the samples with Δ ○ ◎ are 12 times, 24 times, 36 times more excellent than T2, respectively, even when used as a peripheral member of a glass lens mold. Showed life.

  When used as a base electrode used in electric heating forging of various steel parts such as S45C and stainless steel, the circles in the die cast mold row could be used. Compared with T2, ◯ indicates an excellent long life of 10 times or more, and ◎ indicates an excellent life of 20 times or more.

  Since the W-Cr-based alloy of the present invention has excellent oxidation resistance and high-temperature hardness, it can be used for die casting molds, hot extrusion dies, lens mold peripheral members, and electric heating forging electrodes. In order to prolong the service life significantly, it contributes greatly to the industry and is expected to sell over 100 million yen annually.

Claims (8)

In the alloy produced by the powder metallurgy method, the total amount of Ni and Fe in the compounding composition is 1.5 mass% or more and 3.1 mass% or less of the whole alloy, and Cr in the compounding composition is 2.0 mass% or more and 5.0 mass% or less, And the total amount of Cr, Nb, Ta, Ti in the composition is 2.0 mass% or more and 7.6 mass% or less of the whole alloy, the balance is W, and an inevitable impurity is added to this, and the sintered alloy The oxygen content is 0.051 mass% or more and 0.085 mass% or less, and part of the blended W and Cr is a W-Cr solid solution as a dispersed phase, the hardness is 33.8 HRC or more, W-Cr base alloy with a bending force of 830 MPa or more. In the alloy produced by the powder metallurgy method, the total amount of Ni and Fe in the compounding composition is 1.5 mass% or more and 3.1 mass% or less of the whole alloy, and Cr in the compounding composition is 2.3 mass% or more and 5.0 mass% or less, And the total amount of Cr, Nb, Ta, Ti in the composition is 3.0 mass% or more and 7.6 mass% or less of the entire alloy, the balance is W, and inevitable impurities are added to this. The amount of oxygen in the alloy is 0.051 mass% or more and 0.085 mass% or less, and a part of the blended W and Cr is a W-Cr solid solution as a dispersed phase, and the hardness is 41.8 HRC or more. W-Cr base alloy with a bending strength of 830 MPa or more. In alloys prepared by powder metallurgy, the total amount of Ni and Fe in the compounding composition is 1.6 mass% or more and 3.1 mass% or less of the whole alloy, and Cr is 2.3 mass% or more and 5.0 mass% or less in the compounding composition. In addition, the total amount of Cr, Nb, Ta, and Ti in the composition is 4.0 mass% to 7.6 mass% of the entire alloy, the balance is W, and inevitable impurities are added to the composition. The amount of oxygen in the alloy is 0.051 mass% or more and 0.085 mass% or less, and a part of the blended W and Cr is a W-Cr solid solution as a dispersed phase, and the hardness is 43.0 HRC or more. A W-Cr-based alloy having a bending strength of 1572 MPa or more.   A die-casting die for an aluminum alloy, produced using the W-Cr-based alloy according to any one of claims 1 to 3.   A peripheral member of a glass lens molding die produced using the W-Cr-based alloy according to any one of claims 1 to 3.   A platform electrode for electric heating forging of steel parts produced using the W-Cr-based alloy according to any one of claims 2 and 3.   A hot extrusion die for forming a brass bar material produced using the W-Cr-based alloy according to claim 3. For W-Cr-based alloy of any composition of the claims 1 to 3, by the Ar atmosphere at a pressure of atmosphere or 60kPa comprises 600ppm to less _{O 2} or 400 ppm 100 kPa less during sintering, The method for producing a WC-Cr-based alloy according to any one of claims 1 to 3, wherein the amount of oxygen in the sintered alloy is 0.051 mass% or more and 0.085 mass% or less.