EP0362351B1 - Process for producing an ODS sintered alloy and alloy obtainable by this process - Google Patents

Abstract

Process for producing a ductile, high strength, oxide dispersion hardened sintered alloy based on a metal having a high melting point. In the past, oxide dispersion has played only a minor role in comparison with other known processes for increasing strength. The process disclosed permits cost-effective production of metallic materials which possess a strength hitherto unttainable by oxide dispersion and a higher ductility than prior art materials. As a result, the metallic and nonmetallic foreign components in the sintered alloy can be restricted to the relatively small quantities of dispersoids and any dissolved residual oxygen. The process consists in an annealing treatment and calls for a specific choice of basis metal and suitable oxide dispersoid. The materials so obtained are used mainly where metallic components possessing high strength and ductility together with a minimal concentration of foreign elements are required, for example in human medicine where stringent requirements concerning corrosion resistance and biocompatibility apply or in nuclear technology to prevent undesirabble particle reactions.

Description

The invention relates to a method for producing a ductile, high-strength oxide dispersion-hardened sintered alloy from a base metal with a high melting point (T _M ), without fractions, or optionally with fractions of substitution mixed crystal phase that do not have a lasting effect on the alloy properties, in which a metal oxide powder as dispersoid is used to powder the base metal is mixed in, using oxides of metals which have greater formation energies than the oxides of the base metal at temperatures <0.5 T _M.

Classic methods for influencing the strength properties of metals are alloying via mixed crystal phases and mechanical forming. In addition, it is known to increase the strength of materials produced by melt metallurgy or sinter metallurgy by incorporating or separating dispersoids. By definition, dispersoids are understood to mean those particles which are generally built into the metallic base matrix continuously and which do not react with the base metal or dissolve at higher temperatures and are not built into the base lattice as a substitute metal. Oxides, carbides and nitrides are used in particular as dispersoids.

According to firmly established doctrine, the disadvantage of dispersion hardening compared to alloy hardening by continuously or discontinuously separating a second phase in the basic phase from a common solution (precipitation hardening) is that "it is hardly possible to achieve such a degree of dispersity and the increase in strength as they are reached in many cases during excretion processes "(H.Böhm, Introduction to Metallurgy, University Pocket Books, Verlag, Mannheim, Zurich).

To produce dispersion-hardened alloys by powder metallurgical processes, the dispersoids are usually introduced into the powder of the base metal by impregnating the powder with a dispersoid suspension or else by mixing powdered dispersoids.

DE-AS 12 32 353 describes a process for the production of refractory metals, preferably containing metal oxides, by sintered metallurgical processes. The invention there begins with the knowledge that oxides introduced into the powder as colloidal solutions coagulate during sintering, that is to say from the originally fine distribution they form larger oxide particles and thus represent the cause of the high sensitivity to breakage of the sintered material, in particular during further processing . The process feature essential to the invention according to DE-AS 12 32 353 is now based on lowering the sintering temperature compared to the usual standard, in order to reduce the mobility of the introduced oxide particles by diffusion by introducing sintering aids in the form of "other non-harmful metal additives" to the sintered good, which melt at lower temperatures than the oxide and thus enable a kind of liquid phase sintering ("liquid or pasty phase already at temperatures at which no sintering occurs yet. This means that the sintering has ended before the oxide particles - due to lack of movement in the sintered body - become closed have clustered harmful particles ", column 1, lines 51-52, column 2, lines 16-20).
The oxide particles are largely held where they were introduced in the form of the colloidal solution.

Dispersoids introduced in this way can be further homogenized by "mechanical alloying". The aim of mechanical alloying is to distribute the dispersoids as homogeneously as possible even within the individual metal powder grains. These processes are very time-consuming and require high-quality grinding units. They are therefore very expensive and their applicability depends on the nature of the components. Furthermore, in practice a compromise with regard to the degree of homogenization and grinding costs is required; d. H. the grinding is limited in time.

According to DE-A1 35 40 255 for the production of an ODS alloy it is proposed to mix the base metal in the form of a salt solution with the dispersion particles in a colloidal suspension and ultimately to reduce it to metal.
The finely divided, homogeneous introduction of the dispersoid into the metal matrix is mentioned as a particular advantage. But even with this method, the distribution is limited by the particle size of the components.

The production of dispersion-hardened alloys consists, by definition, of incorporating particles which do not react or alloy with the basic matrix as dispersoids. In connection with this, dispersoids with a melting point that is generally well above the alloy sintering temperature are generally used in the sintered metallurgical processes for producing dispersion alloys. The dispersoids are in a solid phase during the entire manufacturing process.

Due to the above-mentioned doctrine that only a comparatively small increase in strength can be achieved by dispersion hardening alone, in the past the means of mixed crystal alloy hardening was additionally used in those cases or precipitation hardening, in which higher strengths were required. In addition to the dispersoids, larger proportions of additional metals were added to the base metals.

In addition to the sintered metallurgical production, it is known to produce oxide dispersion alloys of metals with high melting points by melt metallurgy, in particular by melting in an arc.

From DE-C1 12 90 727 it is known, for example, to produce a high-strength niobium alloy by melt metallurgy by adding small amounts of oxygen, carbon and / or nitrogen to the niobium in addition to 0.5 - 12% zirconium and possibly larger amounts of others metals high melting point is added and the so molten (arc) alloy is between 5 minutes and 9 hours at 1600 to 2100 ^o C solution annealed, cooled, shaped and is finally subjected to a process of Ausscheidungsglühung or paging. In the patent description it is stated that during the solution annealing the second phase, meaning the carbides, nitrides and / or oxides contained in the base matrix (base metal) after melting, forms a solution with the base matrix. According to that invention, by solution annealing and subsequent quenching, the second phase should remain in solution during cold working and should only be excreted homogeneously and finely by aging annealing. The achievable quality is documented by examples and in the form of strength properties compiled in tables.
According to this patent, the means of substitution mixed crystal alloying and the precipitation hardening, according to column 1, line 65 of the description, are used together with the means of dispersion hardening to increase the mechanical properties of such alloys. The strengths achieved thus result from two or three processes running side by side that increase strength and hardness.

The comparatively small amounts of O, but also N and / or C in the alloy prove that there the oxide excretion plays a comparatively minor role as a means of increasing strength. In example 1, the casting block is remelted six times in order to ensure that the metals and dispersoids are homogenized in a manner which is usable, but in no way good due to the process. The process is nevertheless comparatively expensive. After melting and even after hot forming, such alloys have comparatively coarse grains which impair the material strength. For this reason, the description, column 1, line 55 et seq. Also expressly advises "not to prolong the solution annealing of the sheets unnecessarily in order to prevent grain growth".
The room temperature data are not listed in the description. Experience has shown that alloys produced using this process can have relatively high strengths, but at the same time only low ductility at room temperature (see e.g. BVG Grigorovich and EN Sheftel 'Met. Sci. And Heat Treatment 24 (7-8), p. 472, (1983)

US Pat. No. 3,181,946 describes a niobium-based alloy which contains 0.25-0.5% oxygen and / or 1-3% zirconium and / or titanium, the weight ratio oxygen to titanium or zirconium being 3: 1 up to 12: 1 is enough. Material hardening is achieved there by oxide dispersion hardening and, according to the examples given, to a certain extent also by interstitially dissolved oxygen and by alloying niobium with titanium and / or zirconium. It is pointed out there that higher proportions of interstitially dissolved oxygen in the niobium cause great brittleness. In order to counteract this effect, metal oxides of metals with greater binding energy (negative enthalpy of formation) than that of the base metal are only used together with an excess of oxide metal. The patent description is aimed exclusively at the melt metallurgical production of this alloy. The additives are in the arc during a remelting process of the high-purity niobium eg added as titanium oxide powder and sponge-like titanium metal. The cooling process, which is important for the form of dispersion precipitation, is not given any attention in the patent description. This process does not allow a very fine distribution of the dispersoids in the base metal.
The achievable strength properties of additionally hot-formed but not recrystallized niobium alloys are summarized in a table and compared with the properties of pure commercially traded niobium grades. They later serve as a comparison for the strength increases achievable according to the present invention.

The object of the present invention is to develop a more economical process for the production of ODS sintered alloys with high ductility and strength properties using a base metal with a high melting point. The strength values of alloys produced by known metallurgical processes, both in the deformed and in the recrystallized state, should at least be achieved without substitution mixed crystal formation and classic precipitation of a second metal or compound phase being used as a means of increasing the strength.
The process should be able to control the degree of dispersion hardening very precisely. The ductility of the alloy should also be sufficiently large for subsequent cold forming of the material.
The properties of a single metallic element, such as its corrosion behavior and the radiation physical properties, should be preserved as far as possible unaffected by foreign elements and at the same time the mechanical strength of the metals compared to the high-purity phase, with or without deformation hardening, should be significantly increased.

• das eingebrachte Oxid zersetzt sich und/oder wird vom Grundmetall reduziert, die entstehenden Komponenten gehen im Grundmetall in Lösung
• die gelösten Komponenten werden infolge Diffusion im Grundmetall fein verteilt
• ein Teil des insgesamt in der Legierung befindlichen Sauerstoffes dampft von der Oberfläche des Sinterkörpers ab,

und daß die Oxid-Dispersoide während des Abkühlens der Sinterlegierung am Ende des Sintervorganges, oder bei einer anschließenden Auslagerungsglühung, kontrolliert aus der Lösung ausgeschieden werden.This object is achieved according to the invention by a method in which a powder compact formed from the powder mixture is sintered during the sintering process at temperatures in the range 0.7-0.9 T _M while the following processes are taking place

• the introduced oxide decomposes and / or is reduced by the base metal, the resulting components dissolve in the base metal
• the dissolved components are finely distributed due to diffusion in the base metal
• part of the total oxygen in the alloy evaporates from the surface of the sintered body,

and that the oxide dispersoids are eliminated from the solution in a controlled manner during the cooling of the sintered alloy at the end of the sintering process, or during a subsequent aging annealing.

Preferred embodiments of the method are disclosed in claims 2-3. Oxide dispersion-hardened sintered alloys, which can be produced by one of the processes mentioned, are disclosed in claims 4-6.

Due to the features of the invention listed, the method can only be applied to a limited number of alloys. Among the metals with high melting points are primarily those of the V. and VI. Subgroup of the periodic table suitable. Due to the free negative energy of formation, only a limited number of oxides can be used for the desired dispersion hardening. The following table provides an overview of oxides that can be used at least in individual cases and their free formation energy and, in comparison, the oxides of some high-melting metals with comparatively low formation energies: table Solution metal Solution metal oxide and hard temperature-resistant metal oxide Approximate values of the negative free oxide formation energy at 25 ^o C in kilojoules per gram atom of oxygen silicon SiO₂ 403 kJ titanium TiO₂ 424 kJ Zircon ZrO₂ 512 kJ aluminum Al₂O₃ 529 kJ beryllium BeO 584 kJ Thorium ThO₂ 613 kJ chrome Cr₂O₃ 348 kJ magnesium MgO 572 kJ manganese MnO 365 kJ MnO₂ 233 kJ Lanthanum La₂O₃ 580 kJ hafnium HfO₂ 566 kJ barium BaO 529 kJ strontium SrO 560 kJ calcium CaO 605 kJ yttrium Y₂O₃ 604 kJ niobium Nb₂O₅ 357 kJ Tantalum Ta₂O₅ 388 kJ Vanadium VO 416 kJ molybdenum MoO₂ 251 kJ MoO₃ 227 kJ tungsten WO₂ 251 kJ WO₃ 247 kJ rhenium ReO₃ 189 kJ

An important factor determining the respective choice of suitable combinations of base metal and dispersoid is the solubility of the oxygen and the oxide metal in the base metal at the respective sintering temperature and the melting temperature of the oxide metal itself. Insufficient solubility or the formation of intermetallic connections between oxide metal and base metal exclude some combinations from metal and oxide or at least limit the achievable dispersoid content in the alloy.

Each of the three processes that run side by side in solution annealing according to the invention is known per se, as are the means and measures to ensure a controlled process of these processes. It is therefore within the scope of the average specialist to take suitable measures for the desired coordination of the three processes in each case.

The concentration of the oxide in the base metal essentially determines the respective temperature at which the individual processes according to the invention take place or become dominant over the others. By matching the sintering time and sintering temperature to the components present in the respective alloy and their concentration, the three processes for oxide homogenization can be achieved in the course of the sintering.

The total oxygen content in the sintered material should preferably be set so that just the stoichiometrically necessary amount remains for the formation of the oxide, which is strictly only valid for the sinterling center due to a diffusion-controlled concentration profile. In some cases, the oxygen content will be set lower, ie sub-stoichiometric, in order to prevent the oxide from precipitating too quickly and, as a rule, coarse-grained when cooling after the annealing treatment - at the expense of a slight reduction in strength.

Excess oxygen in the sintered material leads to interstitially dissolved oxygen in addition to completely excreted oxide. An oxygen deficit results in incomplete oxide excretion. In the latter case, the oxide metal remains partially dissolved in the basic matrix and thus acts as a getter for impurities in later processing steps - but also as a mixed crystal component.
Since interstitially dissolved excess oxygen, which is not excreted as an oxide, brings about an additional increase in strength, but at the same time leads to a decrease in ductility, in practice an optimum of the influencing variables which has been adapted to all requirements must be established.

The sintering and annealing process can take place both by direct sintering and by indirect sintering. In the direct sintering process, the sintered material is heated by direct current passage. The necessary water cooling of the connections enables the sintered goods to cool down particularly quickly when the sintering process has ended.
Following the sintering process with solution annealing, depending on the dispersoid and concentration, the re-excretion in the form of the finest, homogeneously distributed oxide particles will already take place during cooling or during a subsequent aging annealing. The rate of cooling plays an important role here, and more so the higher the oxide concentration in the alloy. Directly sintered material can be quenched particularly quickly to low temperatures. By heating the alloy, e.g. B. before the extrusion as the first forming process, the precipitation of the oxide particles is occasionally made possible or completed.

In order to carry out mechanical shaping processes, in particular cold forming by forging, rolling or hammering, the oxide dispersion alloy according to the invention must have sufficient ductility in addition to high strength. It is therefore important to determine the strength properties of the alloy according to the invention by choosing the dispersoid concentration, but especially by the Controlled control of the solution annealing according to the invention so that it can be brought as close as possible to a just tolerable limit value.

According to a preferred embodiment of the invention, the alloy consists of niobium or tantalum as the base metal and, in addition to small amounts of dissolved oxygen, contains 0.2-1.5% by weight oxides using one or more of the metals Ti, Zr, Hf, Ba, Sr, Ca, Y, La.

Particularly outstanding results are achieved with a niobium alloy which contains 0.2-1% by weight of TiO₂, whereby in addition to small amounts of interstitially dissolved oxygen, TiO₂ is present in the basic matrix as a finely divided dispersoid in the basic matrix.
Another preferred niobium alloy contains 0.2-1.5% by weight of ZrO₂.

The unusually high strengths achieved by means of the invention combined with a comparatively high ductility for dispersion sintered alloys were surprising and could not have been foreseen to this extent. For example, in the publication "Niobium, TMS-AIME, Proceedings of the International Symposium 1981, ed. H. Stuart (1984)", on p. 247, it is stated that in niobium, due to the lack of sufficiently finely divided dispersoids, very little dispersion solidification can be achieved . The results of the present invention could not even be achieved even in the cases in which approximately comparable annealing processes were used in the case of melt metallurgical alloy formation according to the prior art with regard to temperature and time. Rather, it must be assumed that, because of the different boundary conditions, the three processes which take place in the present method in addition to sintering cannot be coordinated with one another in a comparable manner. In particular, in contrast to sintered alloys, the metal component of the decomposing oxide should evaporate out of the alloy much more easily when the alloy material is melted. It therefore has no possibility of being distributed comparatively homogeneously in the base metal.

To the extent that oxide dispersion alloys have previously been produced by sintering, the sintering has been carried out at comparatively substantially lower temperatures than according to the present invention. This ensured that the oxide particles distributed in the compact remain as unchanged and stationary as possible at their place of introduction.
The feasibility of the annealing treatment according to the invention to the extent that was actually possible was surprising. According to the prevailing doctrine, there was a fear that at the annealing or sintering temperatures according to the invention, in addition to the oxides of the base metal, the dissolved oxide metals would also evaporate at high rates from the sintered body surface. Because taking into account the boundary conditions to be fulfilled for the oxide formation energies, the melting points of the oxide metals can be significantly below the annealing temperatures which are respectively desirable according to the present invention and, in preferred embodiments, are also below the annealing temperatures.

A major advantage of the method according to the invention is its economy. Insofar as dispersion alloys previously produced by melt metallurgy were produced using roughly comparable annealing processes, the entire production process was significantly more cost-intensive with significantly lower solidifications. B. melting and repeated remelting of the oxides in the casting by means of arc melting.

On the basis of the strengths that can be achieved in each case according to the various processes, it can be assumed that for ODS sintered alloys according to the invention, much finer oxide particles and homogeneous dispersoid distributions can be achieved in the basic matrix than with conventional melt metallurgical processes including an annealing treatment. This has the advantage that a much finer grain is regularly obtained during sintering than by melt metallurgy.

A significant economic advantage of the method according to the invention is based on the integration of the annealing treatment according to the invention in the generally required sintering process.

• die ODS-Sinterlegierungen besitzen eine vergleichsweise hohe Duktilität und lassen sich daher wesentlich wirtschaftlicher auf höhere Endfestigkeiten umformen
• die Legierungen sind regelmäßig korrosionsfester als nach bekannten Verfahren hergestellte
• typische für ihre Verwendbarkeit ausschlaggebende Eigenschaften einzelner Grundmetalle, wie extreme Korrosionsbeständigkeit und damit Körperverträglichkeit als Humanimplantat, aber auch die Verwendung beispielsweise des Niobs aufgrund seines niedrigen Neutronen-Einfangquerschnitt, werden durch die geringen Dispersoid-Konzentrationen praktisch nicht beeinflußt.

Comparably high-strength alloys - especially at room temperature and medium-high temperatures - have so far only been obtained for the respective base metal by means of mixed crystal phase formation, possibly with the separation of a second metal phase. A deliberate avoidance of mixed crystal phase formation has the following advantages:

• The ODS sintered alloys have a comparatively high ductility and can therefore be formed much more economically to higher ultimate strengths
• the alloys are generally more corrosion-resistant than those produced by known processes
• Typical properties of individual base metals which are decisive for their usability, such as extreme corrosion resistance and thus body tolerance as a human implant, but also the use of niobium, for example, due to its low neutron capture cross section, are practically not influenced by the low dispersoid concentrations.

Materials manufactured according to the present invention are required in chemistry as well as in tools for the high-performance forming of special alloys, such as super alloys.
An important field of application for niobium and tantalum alloys is implants for human medicine. The use of such high-purity niobium and tantalum alloys, which are known to be particularly tissue-compatible, has so far failed due to their inadequate strength properties. Niobium and tantalum alloys produced by the method according to the invention therefore considerably expand the field of application in implant medicine.

A promising application of alloys according to the inventive method is in pipe systems for alkali metal cooling circuits, for. B. in nuclear power plants.

The outstanding strength properties of alloys according to the invention are shown in connection with the following examples.

example 1

An alloy niobium - 0.5 wt.% TiO₂ is produced according to the inventive method. For this purpose, 3980 g of niobium powder with an average particle size of 10 µm and an oxygen content of <1000 ppm are mixed homogeneously with 20g TiO₂ powder agglomerate with an average particle size of 0.25 µm for 1 hour.
The powder mixture is then hydrostatically compressed to 80% of the theoretical density at approx. 2000 bar.
The compact thus obtained is slowly heated in a high vacuum (better 1 x 10⁻⁵ mbar) and finally sintered at a temperature of 2100 ^o C for 12 hours. These sintering conditions are matched to the size of the samples and the diffusion and degassing processes to be achieved. This leads to the decomposition and solids solution of the TiO₂ and the diffusion of the Ti and O₂ components in the niobium. In addition, a part of the oxygen, especially in the form of niobium oxide, is evaporated from the surface of the sintered body.
The result is a very homogeneous distribution of titanium and oxygen, in a stoichiometric ratio in the core area of the sample and in a slightly sub-stoichiometric ratio in terms of oxygen in the edge area of the sample. Furthermore, it was found that the concentration of the titanium is approximately constant over the entire cross section of the sintered body except for an edge zone in the mm range.
Due to the low TiO₂ concentration in the alloy, no significant TiO₂ excretion takes place during cooling after the sintering process, but an almost complete excretion during an approximately one-hour period Heating and aging process at the beginning of the hot forming process. Electron microscopic examinations of samples after aging annealing showed that the alloy had very homogeneously distributed, fine-grained TiO₂ particles with a particle size of 2 - 20 nm, preferably 8 - 12 nm.
Alloys of this type can be processed further using the known hot and cold forming processes. In the present case, hot forming is first carried out by extrusion at 1000 ^° C. with a forming ratio of 8.7: 1. The alloy sample was then further processed by profile rolling and rotary hammering up to a degree of cold deformation of 72%. The cold forming could easily be increased to 99.9% without intermediate annealing.

Strength tests were then carried out on standard samples made from rods with an 8 mm diameter. In Table 1, the strength values achieved are summarized under item 1. The table shows two sets of tensile strengths at room temperature 800 ^o C, 1000 ^o C and 1200 ^o C, specifically for the deformed sample and for the sample subsequently recrystallized at 1400 ^o C for 1 hour. In addition to the tensile strengths, the table also contains associated elongation values.

In addition to the tensile strength, the fatigue strength of such alloys was also examined. The tests showed above average values at 2.10⁸ cycles using an ultrasound method with approx. 400 N / mm² fatigue strength in air.

The alloy has excellent ductility. This manifests itself on the one hand in the good processability and also in a very low value of approx. -50 ^o C for the transition temperature, in a high notched impact strength of approx. 135 J / cm² at room temperature and in a high elongation at break of> 10% deformed material.

Example 2

An oxide dispersion-strengthened niobium-1 TiO₂ alloy was produced by the process described in Example 1. For this purpose, the TiO₂ was weighed in twice as high as in Example 1.

In contrast to Example 1, a partial excretion of the TiO₂ could already be observed in this case during the cooling after the sintering and reaction annealing process. When the alloy was warmed up before the subsequent hot forming, the titanium still in solution was almost completely excreted as TiO₂.
The higher TiO₂ content in the alloy caused a higher resistance to deformation, so that the samples were advantageously annealed before the individual steps of cold forming in order to achieve a more uniform structure.

The tensile strength and elongation at room temperature measured on this sample are listed in Table 1 under item 2.

Example 3

After the process steps listed in Example 1, a niobium-0.5 ZrO₂ alloy was produced.

The powder compact was further processed by means of direct sintering, especially with a view to rapid cooling of the sintered material after the sintering and annealing process.
Since ZrO₂ is more resistant than TiO₂, the sintering temperature was increased to 2300 ^o C to ensure on the one hand that the components of the ZrO₂ go completely into solution, but on the other hand to set the total oxygen content of the sample somewhat lower and thus too fast and comparatively to prevent rough re-excretion of the oxide when the sample cools after the sintering process. Measures known per se ensured rapid cooling of the sintered product.
Taking into account the higher ZrO₂ stability compared to TiO₂, the heating or aging temperature was increased by 100 ^o C to 1100 ^o C before the first hot forming process.
The further process steps were carried out in accordance with Example 1.
The tensile strengths and elongation values obtained in the formed as well as in the recrystallized state are listed in Table 1 under item 4.

Position 1

eine Nb-TiO₂-Legierung entsprechend Beispiel 1 vorliegender Erfindung

Position 2

eine Nb-TiO₂-Legierung entsprechend Beispiel 2 vorliegender Erfindung

Position 3

eine Nb-1.5 Ti-0.5 0-Legierung entsprechend der zum Stand der Technik zitierten US PS 3 181 945

Position 4

eine Nb-ZrO₂-Legierung nach Beispiel 3

Position 5

eine Nb-1 Zr-Legierung nach dem Stand der Technik ("Niobium, TMS-AIME Proceedings of the International Symposium 1981)

Position 6

eine Niob-1 Zr 0,25 O-Legierung entsprechend der zum Stand der Technik zitierten US PS 3 181 945

Position 7

ein hochreiner Niob-Werkstoff entsprechend Literaturwerten und eigenen Messungen.

Table 1 shows items 1 to 7 tensile strengths and associated elongation values at different temperatures for a number of different samples.
It is

position 1

an Nb-TiO₂ alloy according to Example 1 of the present invention

Position 2

an Nb-TiO₂ alloy according to Example 2 of the present invention

Position 3

an Nb-1.5 Ti-0.5 0 alloy in accordance with US Pat. No. 3,181,945 cited in the prior art

Position 4

an Nb-ZrO₂ alloy according to Example 3

Position 5

an Nb-1 Zr alloy according to the prior art ("Niobium, TMS-AIME Proceedings of the International Symposium 1981)

Position 6

a niobium-1 Zr 0.25 O alloy in accordance with US Pat. No. 3,181,945 cited in the prior art

Position 7

a high-purity niobium material according to literature values and own measurements.

The results according to the invention can only be compared to a limited extent with the cited literature values, because on the one hand the deformation process of the samples according to the cited prior art is not described in detail and because On the other hand, based on the description details given there, it can be assumed that in addition to the oxide dispersion precipitates, the alloy also contains significant proportions of oxide metals of the dispersion oxides in the basic matrix and has an alloy effect which increases strength.
From a purely qualitative point of view, however, it can be stated that, according to the prior art, it is not possible to achieve high strength values comparable with the present invention. The information for pure niobium under item 7 shows that, for dispersion alloys produced according to this invention, significantly higher strengths can be achieved, at least at room temperature, than by means of shaping and possibly recrystallizing pure niobium. TABLE 1 Pos. Material data in% by weight Status Test temp. ^o C Tensile strength MPa Strain % 1 Nb-0.5 TiO₂ deformed RT 950 12th 800 405 12th 1000 350 15 1200 250 18th recist. RT 490 34 800 175 33 1000 135 46 2nd Nb-1 TiO₂ deformed RT 1100 12th recist. RT 535 29 3rd Nb-1.5 Ti- 0.5 0 (U.S. Patent 3,181,946) hot deformed RT 506 29 871 307 19th 982 251 14 1204 185 20th 4th Nb-0.5 ZrO₂ deformed RT 760 11 recist. RT 450 32 5 Nb-1 Zr (Niobium TMS-AIME) deformed RT 350-550 5-15 recist. RT 290 35 800 190 18th 1000 135 32 1200 90 77 6 Nb-1 Zr 0.25 0 (U.S. Patent 3,181,946) hot deformed RT 530 16 982 312 17th 1093 224 26 7 Niobium; purely deformed RT 300-550 2-15 recist. RT 200-300 20-45 8th Ta-0.5 TiO₂ deformed RT 890 13 recist. RT 470 31 9 Ta; purely deformed RT 450-650 2-7 recist. RT 300-350 35-55

Example 4

Analogously to the exemplary embodiments 1-3, an alloy tantalum - 0.5% by weight TiO 2 is produced, the higher melting point of the tantalum having to be taken into account in some process parameters.

7760 g of tantalum powder with an average grain size of 9.5 µm and an oxygen content of 1050 ppm are homogeneously mixed with 39 g of TiO₂ with an average grain size of 0.25 µm (identical oxide powder as in Examples 1-3).

To avoid an excessive decrease in O₂ due to the evaporation of tantalum suboxides (TaO, TaO₂), the sintering temperature is 2300 ^o C compared to the usual approx. 2600 ^o C. This achieves an approximately stoichiometric oxygen concentration corresponding to the titanium concentration introduced. The lower sintering density due to the lower sintering temperature is completely sufficient for complete compression during the subsequent extrusion. The aging annealing for the separation of the finest TiO₂ particles takes place here preferably at 1100 ^o C.
Due to the high heat resistance of tantalum, extrusion is carried out at 1200 ^o C.
Cold forming is then carried out using profile rollers and rotary hammers - a total of approx. 80% forming.

Table 1 shows in item 8 the tensile strengths and elongation values obtained in turn on 8 mm test specimens in the deformed state and after recrystallization. The high recrystallization temperature (1600 ^o C / 1 h) leads to a clear coarsening of the TiO₂ dispersoids and thus to a weakening of the dispersion hardening compared to the cold-formed material. The combination of work hardening and dispersion hardening results in particularly high strengths while maintaining sufficient ductility.
For comparison, values of pure tantalum with 82% deformation are given in position 9, the production steps and process parameters corresponding to the above.

Claims (6)

1. Procedure for the manufacture of a ductile, high-strength oxide dispersion-hardened sintered alloy from a base metal with a high melting point (T_M), without fractions, or where necessary with substitution mixed crystal phase fractions not adversely affecting the alloy properties, in which a metal oxide powder as dispersoid is added to the base metal powder, whereby oxides of such metals as, at temperatures of < 0.5 T_M, have greater energies of formation than the base metal oxides are used,
   characterised in
   that a powder compact formed from a powder mixture is sintered during the sintering process at temperatures ranging from 0.7-0.9 T_M with progression of the following processes:

   - the introduced oxide decomposes and/or is reduced by the base metal, and the generated components enter into solution in the base metal;

   - the dissolved components are finely distributed in the base metal by diffusion;

   - some of the oxygen altogether present in the alloy evaporates from the surface of the sintered body;

   and that the oxide dispersoids are separated from the solution in a controlled manner during cooling of the sintered alloy at the end of the sintering process or during an ensuing precipitation-annealing process.
2. Procedure according to claim 1, characterised in that the sintered alloy is manufactured by direct sintering of the powder compact.
3. Procedure according to claim 1 or 2, characterised in that some of the oxygen as base metal oxide evaporates from the surface of the sintered body.
4. Oxide dispersion-hardened sintered alloy which can be manufactured by a procedure according to claims 1 to 3, characterised in that this consists of niobium or tantalum base metal and contains, apart from small amounts of dissolved oxygen, 0.2-1.5 wt% oxides with application of one or more of the metals Ti, Zr, Hf, Ba, Sr, Ca, Y, La.
5. Oxide dispersion-hardened sintered alloy according to claim 4, characterised in that this is a niobium alloy containing 0.2-1 wt% TiO₂.
6. Oxide dispersion-hardened sintered alloy according to claim 4, characterised in that this is a niobium alloy containing 0.2-1.5 wt% ZrO₂.