JPH11286770A - High corrosion resistance molybdenum-based composite material and its production - Google Patents

Abstract

PROBLEM TO BE SOLVED: To impart a corrosion performance equal to that of Ta to a material and to obtain a mechanical strength and hardness more excellent than those of Ta by subjecting Mo and an Mo-based alloy to nitriding treatment. SOLUTION: This high corrosion resistance Mo-based composite material is characterized by providing the surface of an Mo alloy with an Mo2 N layer of 0.5 to 10 μm thickness. The method for producing a high corrosion resistance Mo-based composite material is characterized by subjecting an Mo series alloy to nitriding treatment for 0.2 to 100 hr in an atmosphere heated at a temp. of 700 to 1,150 deg.C in the presence of gaseous N2 or gaseous NH3 .

Description

DETAILED DESCRIPTION OF THE INVENTION

[0001]

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to a Mo alloy having a high strength and a high corrosion resistance equal to or higher than that of Ta.
_{The present} invention relates to a high corrosion-resistant Mo-based composite material provided with a 2N layer and a method for producing the same.

[0002]

2. Description of the Related Art Conventionally, as a highly corrosion-resistant material under severe conditions, a stainless steel container with a Ta lining is known.

The Mo metal used in the present invention includes γ-Mo ₂ N (face-centered cubic, fcc) and β-Mo ₂ as nitrides.
N (body-centered square, considered bct), δ-Mo
There are three phases, N (hexagonal hcp and non-equilibrium phase BI form). However, there are few studies on Mo nitrides since these nitrides hardly form in pure N ₂ gas which is easy to handle. The β phase is a phase that is stable only in a low temperature range of 700 ° C. or lower, and has a C-axis length twice that of the γ phase.
ct structure. If the amount of nitrogen is small, the γ phase can be obtained by rapid cooling from 700 ° C. or higher even in the composition range. On the other hand, δ-MoN is known to be synthesized by nitriding MoO ₃ or MoS ₂ with NH ₃ gas. The crystal structure of δ-MoN has a lattice constant a = 0.5
725 μm, c = 0.5608 μm hexagonal system (hc
p) Structure, but details are unknown. As described above, the details of Mo nitride, such as its physical properties, crystal structure, and formation temperature range, are unknown, and the high corrosion resistance of the treated product is not known at all.

[0004]

The Ta used in the present invention has excellent performance in terms of corrosion resistance, but has low mechanical strength, is significantly worn or deteriorated, and is difficult to use for a long time (has no durability). ), High specific gravity, and high cost.

Although it has been known that surface nitriding of a Mo-based alloy can be performed, it is not known how excellent it is, what characteristics it has, and how it can be used. was there.

[0006]

SUMMARY OF THE INVENTION The present invention provides an excellent corrosion resistance and a high mechanical strength by subjecting a Mo alloy to a nitriding treatment under a certain condition and providing a thin layer of Mo nitride on the surface thereof. And succeeded in obtaining an inexpensive and lightweight composite material.

Namely invention materials are high corrosion resistance Mo-based composite material characterized in that a Mo ₂ N layer having a thickness of 0.5μm~10μm the surface of the Mo alloy, Mo ₂ N layer,
This is a β-Mo ₂ N layer.

Further, the invention of the method is characterized in that the Mo-based alloy is subjected to nitriding treatment in an atmosphere of 700 ° C. to 1150 ° C. for 0.2 hours to 100 hours in the presence of N ₂ gas or NH ₃ gas. This is a method for producing a Mo-based composite material. Next, the invention of another method is that a Mo-based alloy is nitrided in an atmosphere of 700 ° C. to 1150 ° C. in the presence of N ₂ gas or NH ₃ gas and has a thickness of 0.5 μm or more that does not cause cracks. A high corrosion-resistant Mo-based composite material characterized by providing a Mo ₂ N layer of
It should be less than 0 μm.

In the above invention, the nitriding temperature is 700
If the temperature is lower than 1 ° C. or higher than 1150 ° C., a desired Mo ₂ N layer having excellent corrosion resistance cannot be obtained. Also, when the thickness of the Mo ₂ N layer is less than 0.5 μm or more than 10 μm, it is difficult to obtain the corrosion resistance aimed at by the present invention. In the case of the β-Mo ₂ N layer, excellent corrosion resistance is certainly exhibited.

In the above, when the ambient temperature exceeds 1200 ° C., no Mo ₂ N layer is formed, but this is presumed to be due to decomposition. If the ambient temperature is lower than 700 ° C., the growth of the Mo ₂ N layer is remarkably slowed down, which is not suitable for industrial production. For example, it has been found that it takes 100 hours or more to perform a nitriding treatment at 700 ° C. to obtain a required thickness (1 μm or more).

In the present invention, when the processing temperature is high, M
The time for growing the o ₂ N layer to a predetermined thickness is reduced,
Low temperatures take a long time. For example, 800 ° C., 1
In time, a 1 μm thick Mo ₂ N layer can be obtained,
At 1000 ° C. for 1 hour, a Mo ₂ N layer having a thickness of 12 μm was formed.

Therefore, the thickness of the Mo ₂ N layer can be controlled by controlling the nitriding temperature and time.

[0013]

BEST MODE FOR CARRYING OUT THE INVENTION According to the present invention, a Mo-based alloy is subjected to a nitriding treatment so that the Mo-based alloy has a thickness of 0.5
Mo system provided Mo ₂ N layer of μm~10μm a composite material or a manufacturing method thereof.

The Mo ₂ N layer is preferably β-Mo ₂ N. In order to grow the thickness to 0.5 μm to 10 μm, the Mo ₂ N layer should be in an atmosphere at 700 ° C. to 1150 ° C. for 0.2 hours to 1 hour.
Process for 00 hours.

(Experimental Example 1) High-purity Mo powder and Ti
A compact was prepared using C powder as a raw material,
The sintered body was obtained by sintering in a hydrogen atmosphere at ℃. This sintered body is hot- and warm-rolled, and further cold-rolled to a thickness of 1 mm.
Was obtained. From this plate material, a square rod-shaped sample (2 mm ^W x 1
cut out mm ^E × 25mm ^L), after polishing the surface by emery paper, it was subjected to electrolytic polishing. Then, in a vacuum (about 1.
It was recrystallized by a heat treatment at 1500 ° C. for 1 hour at 3 × 10 ⁻⁴ Pa) to obtain a sample for nitriding. Nitriding is 1 atm NH
_This was performed for 4 to 16 hours in a stream of _three gases.

The cross section of the sample obtained above was mechanically polished, and then the structure was observed with an optical microscope and a hardness test was performed with a Vickers hardness tester (load: 25 gf, holding period: 15 s). The etching for the optical microscope observation is performed by using Murakami reagent (K ₃ [(Fe (CN) ₆ ]: NaO
H: H ₂ O = 1: 1: 10 (weight ratio). In order to identify the nitride layer on the sample surface and inside, X-ray diffraction (XRD, Cu-Ka, Rigaku Geige) was used.
r Flex). Further, the fine structure of the nitride layer was examined by a transmission electron microscope (TEM, Topcon EM-0).
02B) Observation and examination by electron beam diffraction (ED). The sample for electron microscope is about 1
A thin plate having a thickness of 50 μm was formed, and a sample was formed using a dimple at the center of the sample, and then manufactured using an ion milling device. For ion milling, Ar ⁺ ions are 5 to 7k
V, accelerated at 2 mA and performed at an incident angle of about 10 °.

As shown in FIG. 1, the samples obtained by the above-mentioned treatment were cubic (fcc) γ-Mo
Only the diffraction line of ₂ N was observed. This result indicates that a relatively thick γ-Mo ₂ N is formed on the sample surface by nitriding.

Further, the Mo—Ti alloy is (a) 600 ° C.
(F) The XRD pattern on the sample surface which was subjected to N ₂ gas nitriding at 1100 ° C. for 16 hours was as shown in FIG.

Next, FIGS. 3 (a) and 3 (b) show that 4 ° C. at 1100 ° C.
1 shows optical micrographs of cross sections of pure Mo samples nitrided for 16 hours and 16 hours. In each of the samples, a white layer parallel to the sample surface was formed uniformly, and it can be seen that the thickness of this surface layer increased as the nitriding time was extended. In consideration of the result of FIG. 1B, it is determined that this surface layer is a Mo nitride layer. In addition, this surface layer and the internal Mo
Interface is very smooth. Therefore, it is considered that the grain boundary diffusion of nitrogen hardly contributes to the growth of the surface layer. Hereinafter, this surface layer will be referred to as a surface (Mo) nitride layer.

In general, when such a compound layer grows, there are two types of rate-determining formulas depending on the difference in the rate-controlling mechanism. That is, in the case where the reaction is rate-controlled, the thickness (E) of the compound layer is proportional to the reaction time (t). The following linear law equation (1) holds.

[0021]

(Equation 1) On the other hand, in the case of the diffusion-controlled mechanism, the following parabolic law equation (2) in which E is proportional to the square root of t holds.

[0022]

(Equation 2) Here, _Kp is a growth rate constant. Thus, FIG. 4 shows the relationship between the thickness E of the surface nitride layer obtained by nitriding the pure Mo and Mo-0.5Ti alloys at 1100 ° C. and the square root of the nitriding time t. It was found that a linear relationship was established between the two, and the parabolic law was established. Here pure Mo and Mo-
Each value of _{K p} of 0.5Ti alloy 3.25 × 10 ^-4
mm ′s ^−1/2 , 2.96 × 10 ⁻⁴ mm ′s ^−1/2 . This result suggests that the growth of the surface nitride layer is apparently limited by the diffusion of Mo or nitrogen in the surface nitride layer.

In order to examine whether the growth of the surface nitride layer proceeds outward or inward from the sample surface, a thickness W _b before and after nitriding (here, 1000 _b ) is used.
μm with the thickness of the sample) the and W _a was measured to examine the change with nitridation time of the difference △ W = W _a -W _b. The results are shown in FIG. 5 for the Mo-0.5Ti alloy. 1
As in the case of pure Ni oxidized at 200 ° C. or more, if the growth of the surface nitride layer proceeds by outward growth due to diffusion of Mo to the sample surface, the thickness of the nitrided sample is It is expected to double the thickness of the layer. However, as shown in FIG. 5, the thickness of the sample hardly changed before and after nitriding, and similar results were obtained in the case of pure Mo. Therefore, it is determined that the growth of the surface nitride layer of the present invention is inward growth inward.
As described above, NH of pure Mo and Mo-0.5Ti alloy
It is concluded that the diffusion of nitrogen into the sample is rate-limiting for the growth of the surface nitride layer by ₃ nitriding. This result indicates that the Mo-based material can be nitrided without forming a void due to diffusion of Mo ions to the sample surface.

FIG. 6 shows pure Mo and Mo nitrided for 16 hours.
3 shows a hardness distribution of a cross section of a sample of -0.5Ti alloy. Hv near the outermost surface of the sample is about 1800, and pure Mo before nitriding is used.
It can be seen that the hardness is significantly higher than that of the Mo-0.5Ti alloy (Hv-200). Further, the hardness of the surface nitride layer showed a tendency to decrease toward the inside of the nitride layer. This decrease in hardness was presumed to be due to a phase change in the surface nitride layer. Thus, for the sample nitrided for 16 hours, it was examined how the nitride changed phase from the sample surface toward the inside. Specifically, an operation of performing X-ray diffraction after carefully polishing and removing a predetermined thickness in parallel from the sample surface toward the inside was repeated. The result is shown in FIG. 7 for pure Mo. Depth d = 0
(A) to 25 (b) Cubic (f)
(While a gamma phase), it from the inside (d = 35 (c) cc Structure) γ-Mo ₂ N of to 70 (e) [mu] m) in tetragonal β-Mo ₂ N (β phase) You can see that it is generated. Furthermore, d = 80 μm (f) inside is an unnitrided layer, and only the Mo diffraction peak was observed. It is considered that the phase change of these Mo nitrides causes a difference in hardness within the surface nitride layer. The surface nitride layer of the Mo-0.5Ti alloy was the same as the surface nitride layer of pure Mo.

Therefore, pure Mo will be described below. FIG. 8 shows the change in lattice constant of surface Mo nitride with depth, determined from the result of X-ray diffraction. FIG. 8 shows two interesting results. First, the lattice constant changes discontinuously at a depth of about 28 μm in the surface nitride layer. From this result, it can be inferred that the region where γ-Mo ₂ N on the surface side and β-Mo ₂ N on the inner side coexist hardly exists or is very narrow, if at all. Second, the lattice constants of the γ phase (d <30 μm) near the surface and the β phase (d> 30 μm) inside the nitride layer both change depending on the sample depth.
It can be seen that the a-axis length of the phase and the c-axis length of the β phase are clearly shorter. The amount of nitrogen differs in each case (28.4 at%
〜34.6%) As a result of obtaining lattice constants of γ-Mo ₂ N and β-Mo ₂ N by X-ray diffraction, the lattice constants of these nitrides (a-axis of γ phase, c-axis of β phase) It can be seen that the amount decreases linearly as the amount decreases. Therefore, the phase change of Mo nitride in the surface nitride layer in the present invention is
It is thought to be due to the change in concentration (gradient) of nitrogen in the γ-Mo ₂ N and β-Mo ₂ N phases.

According to FIG. 9, when the Mo nitride in the surface nitride layer changes from the γ phase to the β phase, the structure changes from cubic to tetragonal and the value of c / a deviates from 1. . However, from the results in FIG. 9, the axial ratio of the β phase is 0.966 to 0.95.
3, which can be regarded as a value almost close to 1. That is, it can be said that the β phase is a tetragonal crystal close to a cubic crystal. Further, the c / a of the β phase will be examined by focusing on the lattice of Mo atoms in the crystal structure. FIGS. 10A and 10B are schematic diagrams of the crystal structures of γ-Mo ₂ N and β-Mo ₂ N. The structure of the β phase used the most reliable data from Evans and Jack. The structure of the γ phase can be regarded as a structure in which nitrogen atoms penetrate into lattice positions (octahedral gaps) of the fcc lattice formed by Mo atoms. On the other hand, the structure of the β phase is similar to that of the γ phase, but considering the unit cell of the crystal, the β phase is considered to be fct whose c-axis is slightly (0.018 nm) shorter than the a and b axes. Can be. Also, from FIG. 9, the axis ratio of the β phase is almost 1. Therefore, in the present invention, the β phase is hereinafter assumed to have an approximately fcc structure.

In the β phase, as shown in FIG.
A number of martensitic structures similar to the rapidly quenched structure are observed, and countless stripes are observed in mutually parallel bands having a width of about 0.2 μm. It is considered that in this region, the stress is relaxed due to the difference in the coefficient of thermal expansion between the surface Mo nitride layer and the mother phase Mo, and misfit of the lattice. FIGS. 12A and 12B show a part of the result of electron beam diffraction of the β phase together with a transmission electron microscope photograph. Here, the incident directions of the electron beams are all [100]. FIG. 12
(A) and (b) are diffraction patterns from adjacent band regions (a and b in the figure), respectively. Comparing the two diffraction patterns, it can be seen that adjacent bands have basically the same crystal orientation. Also, band b
A streak suggesting the presence of a plate-like substance in a direction perpendicular to the stripes in the band is observed in the diffraction pattern (b) from FIG. From this result, it is considered that the stripe in the band is a trace of a plane defect perpendicular to the sample [100]. Trace analysis was performed on this fringe from electron diffraction images in various directions using stereographic projections. FIGS. 13A and 13B show a part of the result, in which FIG. 13A shows [100] incidence, and FIG.
Is an analysis diagram in the case of [211] incidence. As a result, the direction of the stripe is [011] in the case of [100] incidence (FIG. 13).
(Indicated by an arrow in the figure), and it was found that the incidence was [1-31] in the case of the [211] incidence. Therefore, the plane index of the stripe is included in the [011] zone and the [1-31] zone, respectively, and the plane index of the stripe is the great circle of the [011] zone and the [1-31] zone shown in FIG. Exists on the great circle of the obi. FIG. 14 shows that the plane index common to both is {011}. That is, in the case of FIG. 12A, the (0-11) plane,
In the case of (b), it is the (-101) plane. According to the above examination, the stripes innumerably existing in the band in the β phase are {100}.
It is considered to be a plane defect on the {011} plane perpendicular to the plane.

Next, surface defects on the {011} plane of the β phase will be discussed. As typical surface defects of the inorganic solid material, two types of stacking faults and twins can be considered. The structure of the β phase is basically an fcc structure as shown in FIG. It is necessary to consider a surface defect on the {011} plane for this structure. Considering the atomic arrangement in the plane, the {011} plane is stacked in the form of ABABAB... In which two types of lattice planes are stacked on each other as shown in FIG. Stacking faults in this structure result in AA or BB stacking.
Since this laminated structure generates a very large strain energy, it can be said that the possibility of the {011} plane stacking fault is very low.

Next, the possibility that the plane defect in the β phase is a twin with the {011} plane as a twin plane will be examined. If the structure of the β phase is the same cubic fcc as the γ phase, then no such twins exist. This is because no matter what lattice plane is orthogonal to the {011} plane, the lattices on both sides of the {011} plane are mirror-symmetrical to each other from the beginning. However, in practice, the β phase is a tetragonal crystal whose c-axis is only slightly shorter (c / a = 0.966 to 0.95).
3) A schematic representation of the atomic arrangement of the (011) plane of the β phase is as shown in FIG. In this atomic arrangement, when atoms are slightly displaced in the [0-11] direction on the (011) plane as shown in FIG. 15B, (011) [0-11] which is mirror-symmetric with respect to the (011) plane. ] Type twin (shear strain s = 0.087) can be considered. Since the twin plane is the (011) plane and the plane rotated by shearing is the (0-11) plane, the shear strain s due to the twin is 2φ = 87.5 °, and the equation (2.3) is obtained. From s =
It can be evaluated as 0.087. This value is the above {111} <
It is significantly smaller than s = 0.077 for the 112> type twin. This result indicates that twins of the (011) [0-11] type are likely to be generated. This type of twin is actually observed in In, which is the only fct metal in pure metal. From the above study, it is found that (0)
11) It is considered that [0-11] type twins were generated.
Therefore, an attempt was made to directly observe such an atomic arrangement by TEM. FIG. 16 shows a β-phase high-resolution electron microscope (HR-TEM) image taken at a magnification of 550,000. The sample surface is the same (011) surface as in FIGS. 15A and 15B. From the figure, it can be seen that the lattice fringes on the (020) plane and the (002) plane of the β phase are considerably distorted. However, FIG.
No atomic arrangement as shown in (b) was observed. It was later found that this is because, when an electron beam is converged when observing a high-resolution image, the β-phase structure easily collapses due to damage caused by electron beam irradiation. So the magnification is 83,000
We tried to shoot under conditions that reduced the damage by electron beam irradiation as much as 1000 times. FIG. 17 shows the result.
The sample surface is the same (100) surface as in FIG. [0-1
In the stripes having a width of about 10 nm extending in the [1] direction, the existence of lattice stripes of (002) plane which are almost orthogonal to each other was recognized. FIG. 18 is an inverse Fourier transform image of the enlarged image of FIG. For comparison, (011) in FIG.
A schematic diagram of twins is also shown. The angle formed by the lattice fringes of the (002) planes which are substantially orthogonal to each other and appear in the inverse Fourier transform image is about 87.5 °, and this angle is (002) in the schematic diagram of the (011) twin shown below. It can be seen that the angle substantially matches the angle 88.7 ° (calculated value) formed by the surfaces. From the above results, the countless stripes shown in FIG.
It was determined to be a {011} <011> type twin with the {11} plane being a twin plane. From the above results, in pure metal, I
very rare type {011} <01 which can only be seen in n
It has been clarified that 1> type twins also occur in β-Mo ₂ N nitride.

(Experimental Example 2) The corrosion resistance of the product of the present invention will be described. The sample is pure Mo produced by a powder metallurgy method. A square sample of 10 mm square or a rectangular sample of 2 mm width and 20 mm length was cut out from a plate material having a thickness of 1 mm and 1.5 mm, and an oxide film on the surface was polished and removed with emery paper, and then electropolished. . The sample in the form of a strip is then placed in a vacuum (about 1.3 × 10 ⁻⁴ Pa) to reach 1
It was recrystallized by heat treatment at 500 ° C. for 1 hour. Nitriding was performed at 950 ° C. for 0.5 to 9 hours in an NH ₃ gas stream, and after nitriding, the sample was cooled by moving the sample to a greenhouse portion of a furnace tube.

The corrosion test was carried out using a corrosion test apparatus as shown in FIG. 19, in 40-75 wt% boiling sulfuric acid, or 5-7%.
Performed in 0 wt% nitric acid (normal temperature) for 3 minutes to 8 days. For a corrosion test in boiling sulfuric acid, a 10 mm square sample was used.
For the corrosion test in nitric acid, a strip-shaped sample subjected to a recrystallization treatment was used.

With respect to the obtained sample, the surface state before and after corrosion and the corrosion state of the cross section were measured by a scanning electron microscope (SEM,
(JEOL JSM6300) and an optical microscope. In addition, X-ray diffraction (XRD, Cu-Ka, Rigiaku Geig) was used to identify the sample surface phase before and after corrosion.
er Flex). The erosion degree (corrosion rate) was calculated from the weight change before and after corrosion using the following equation (3).

[0033]

(Equation 3) Here, W ₁ : weight before corrosion (g) W ₂ : weight before corrosion (g) H: corrosion time d: density (g / cm ³ ) s: surface area (cm ² ). As the corrosion resistance criterion, the corrosion criterion for metal materials generally used at present shown in Table 1 was used.

[0034]

[Table 1] However, this criterion is a criterion that focuses on whether it can be mechanically endured in configuring an industrial product production apparatus.

Therefore, when other factors (for example, contamination of a drug by an eluting metal, a short expected use period, etc.) become more important, it is necessary to use different standards for each case. Table 2 shows the relationship between the sulfuric acid concentration and the boiling point.

[0036]

[Table 2] FIG. 20 shows the corrosion rates of various metals in boiling sulfuric acid. Pure Mo exhibits excellent corrosion resistance in boiling sulfuric acid having a relatively low concentration of 60% or less. However, it can be seen that the corrosion resistance is rapidly lost as the concentration exceeds 60% and the oxidizing property of sulfuric acid increases.

FIGS. 21 (a), 21 (b) and 21 (c) show 950 in FIG.
The SEM photograph of the surface before and after the corrosion test in various boiling sulfuric acids of the sample nitrided at 9 ° C. for 9 hours is shown. The test period is 8 days. From FIG. 21, 40% (a), 60% (b), 7%
At any concentration of 5% (c), no change was observed in the state of the sample surface before and after corrosion, indicating that the sample was hardly corroded by boiling sulfuric acid. The result is that NH
Mo nitride layer (Mo ₂ N) formed by ₃ gas nitriding
Suggests that pure Mo exhibits excellent corrosion resistance not only in the low concentration region but also in boiling sulfuric acid in the concentration region of 60% or more where the corrosion resistance is lost.

FIG. 22 shows the corrosion rate of a sample nitrided at 950 ° C. for 9 hours in boiling sulfuric acid. Pure Mo for comparison
The corrosion rate is shown by the solid line. FIG. 22 shows that the Mo nitride layer formed by NH ₃ gas nitriding shows complete corrosion resistance at a sulfuric acid concentration of 60% or less. However, at a sulfuric acid concentration of 75%, no change was observed in the state of the sample surface (FIG. 21 (c)), but about 1/20 of that of pure Mo.
However, it was found that corrosion progressed at a considerable speed (about 0.4 mm / y).

Table 3 shows the sample thickness before and after corrosion of the sample nitrided at 950 ° C. for 9 hours.

[0040]

[Table 3] The test for 8 days at the corrosion rate shown in FIG.
%, The thickness of the sample is calculated to decrease by about 18 μm on both sides. However, from Table 3, it can be seen that the thickness of the sample hardly decreased even after the test in 75% sulfuric acid for 8 days. That is, it was found that corrosion of the Mo nitride layer hardly progressed even in 75% sulfuric acid.

FIGS. 23 (a) and 23 (b) show optical micrographs of the sample surface before corrosion of the sample nitrided at 950 ° C. for 9 hours. FIG. 23 (a) shows a 10 mm square sample, and FIG. This is a strip sample that has been subjected to a recrystallization treatment before nitriding. From these results, it can be seen that both (a) and (b) have cracks on the sample surface after nitriding which are considered to have been caused by strain generated during cooling of the sample (indicated by arrows in the figure). As is clear from FIG. 23B, these cracks propagate not through the grain boundaries of the Mo nitride but through the nitride grains. Therefore, the reason why the corrosion progresses despite the fact that there is no change in the sample thickness before and after the corrosion is that the crack reaches the interface between the nitride layer and the parent phase Mo, and the parent phase Mo is selected at the tip of the crack. This is considered to be due to corrosion.

FIG. 24 shows an optical microscope photograph of the cross section of the sample before corrosion (a) and the sample (b) subjected to the corrosion test in 75% boiling sulfuric acid for 8 days. As indicated by arrows in the figure, the cracks generated in the Mo nitride layer completely penetrate the nitride layer, and it is clearly observed that only the parent phase Mo is selectively corroded at the tip. In addition, comparing FIGS. 24 (a) and 24 (b), even after the test for 8 days, no change was observed in the thickness of the nitride layer (about 30 μm) and the width of the crack before and after the corrosion.
o It can be seen that the nitride layer is hardly corroded (FIG. 24 (b)). The reason why the surface of the nitride layer looks uneven is that the nitride has been chipped off by polishing.

According to the above results, the Mo nitride layer formed by NH ₃ gas nitriding shows extremely good corrosion resistance to boiling sulfuric acid, but the mother phase and sulfuric acid pass through cracks which are considered to be generated by strain generated during cooling of the sample. As a result of the contact of the solution, it was found that in a concentration region higher than 60% where Mo loses corrosion resistance, selective corrosion of Mo progresses at the tip of the crack. Therefore, in order to use the Mo nitride layer as a corrosion resistant material, it is essential to completely suppress the occurrence of cracks.

In general, it is known that when a crystalline substance grows as a film on another crystal, the strain energy increases in proportion to the film thickness. Therefore, it was considered that if the nitriding time is shortened and the thickness of the Mo nitride layer is reduced, the strain energy can be suppressed low, so that the generation of cracks during cooling can be controlled. FIG. 25 shows an optical microscope photograph of the surface of a sample obtained by nitriding the recrystallized material at 950 ° C. for 1 hour (a) and 0.5 hour (b). According to FIG. 24, the thickness of the nitride layer generated by nitriding for 9 hours is about 30 μm. Therefore, N
Regarding growth of Mo nitride layer by H ₃ gas nitriding (2)
Assuming that the parabolic law of the equation holds, the thickness of the nitride layer generated by nitriding at 950 ° C. for 1 hour and 0.5 hour is estimated to be 10 μm and 7 μm, respectively. When the nitriding time is 1 hour (a), cracks are still observed at the portion indicated by the arrow, whereas when 0.5 hours (b), similar cracks are not observed. Therefore, the nitriding time is shortened. It has been clarified that the reduction of the thickness of the Mo nitride layer is extremely effective for controlling crack generation.

FIG. 26 shows the XRD of the surfaces of the nitrided material for 9 hours (a), the nitrided material for one hour (b) and the nitrided material for 0.5 hour (c).
Indicates a pattern. The surface layer of the 9-hour nitride material (a) and the 1-hour nitride material (b) is γ-Mo ₂ N, whereas
In the case of the nitrided material (c) for 5 hours, a slight diffraction line of the γ phase is recognized, but it can be seen that the main phase is the β phase. The result is
This is also a very important result in that it suggests that the nitride formed at the beginning of nitriding is a β phase. The reason why the Mo diffraction line is recognized in FIG. 26C is that the thickness of the Mo nitride layer is small. It is considered that the reason why the crack generation is controlled in the nitrided material for 0.5 hour is that the nitride layer is thin, but the strain may have been relaxed by {011} twins generated in the β phase. This is an issue for future study.

Table 4 shows the corrosion rates of the 9-hour nitriding material and the 0.5-hour nitriding material in 75% boiling sulfuric acid and the sample thickness before and after nitriding. The values of pure Mo are also shown for comparison. FIG. 27 shows a comparison of corrosion rates of the unnitrided Mo material (a) and the nitrided material in NH ₃ gas at 950 ° C. for 9 hours (b) and 0.5 hours (c) in 75% sulfuric acid. As a result of suppressing the occurrence of cracks by shortening the nitriding time and making the nitride layer thinner, the corrosion rate of the nitride material for 0.5 hour is about one third of that of the nitride material for 9 hours (about 1/60 of Mo). ), And a remarkable effect of reducing the thickness of the nitride layer is recognized. However, despite the fact that the sample thickness of the 0.5-hour nitride material did not change at all before and after corrosion, and that the β-Mo ₂ N phase is considered to exhibit excellent corrosion resistance like the γ-phase of the 9-hour nitride material, It can be seen that corrosion is progressing.

FIG. 28 shows an optical microscope photograph of the surface of the nitrided material for 0.5 hour subjected to a corrosion test in 75% boiling sulfuric acid for 5 days. No cracks were observed in the nitrided material for 0.5 hours before corrosion (FIG. 25 (b)), but when the sample after corrosion was carefully observed, it was clear that cracks were generated as indicated by arrows. It became. Considering that the Mo nitride layer is hardly corroded, it is difficult to think that these cracks existed before corrosion, which were so small that they could not be seen by optical microscopy.
Therefore, it is considered that as a result of these cracks occurring slightly during the corrosion test, corrosion progressed apparently at the speed shown in Table 4. It is well known that austenitic stainless steels produce stress corrosion cracking in chlorides. A similar phenomenon may have occurred in the Mo 0.5 hour nitride material of the present invention, but it is necessary to clarify the details by measuring residual stress in the future.

Here, for the purpose of suppressing the generation of cracks, the surface Mo nitride layer was formed in NH ₃ gas at 950 ° C. and 0.5 ° C.
Attempts were made to perform nitriding at 800 ° C. in N ₂ gas with the aim of making it much thinner than in the case of time nitriding and dramatically improving corrosion resistance. As one example, FIG. 29 shows an XRD pattern of the surface of a material nitrided at 800 ° C. for 48 hours in N ₂ gas. The diffraction peak of β-Mo ₂ N is clearly observed in addition to the enhanced diffraction peak of Mo, suggesting the formation of extremely thin β-Mo ₂ N, which is estimated to be about 0.5 to 1 μm. FIG. 27 shows the results of a corrosion test of this nitrided material in a 75% concentrated sulfuric acid solution. According to this figure, the corrosion rate of the nitrided material at 800 ° C. in N ₂ gas is 0.02 mm / y, and the corrosion rate is significantly improved compared to the nitrided material at 950 ° C. for 0.5 hour in NH ₃ gas. It has been found that the composition exhibits complete corrosion resistance even in a sulfuric acid solution at a concentration of 2%. In other words, it has been clarified that by forming an extremely thin Mo nitride layer on the material surface, a Mo-based composite material having surprisingly high corrosion resistance can be produced even in a boiling concentrated sulfuric acid solution.

[0049]

[Table 4] FIG. 30 shows the corrosion rates of various refractory metals in boiling nitric acid. Ta, Nb and Zr show good corrosion resistance at all concentrations, while Ti and W tend to have a slightly higher corrosion rate at intermediate concentrations.

FIG. 31 shows that pure Mo (a) and 950 ° C.
The corrosion rate of the sample (b) in which nitriding was performed for 9 hours in nitric acid at room temperature is shown. Pure Mo (a), unlike the other high melting point metals shown in FIG. 30, has a low passivation ability against nitric acid and is corroded at a considerable rate even at room temperature. If the nitric acid concentration exceeds 10%, it will not show any corrosion resistance,
It shows a maximum around 50% concentration. Pure Mo showed complete corrosion resistance at room temperature at a nitric acid concentration of 5% or less. On the other hand, in the case of the sample (b) in which nitriding was performed for 9 hours, the tendency was similar to that of pure Mo, but the corrosion rate (about 0.33 mm / y to 240 mm / y) in the concentration region of 10% to 50% was It turns out that it is smaller than pure Mo (about 3 mm / 10 ⁴ mm / y) by one digit or more. The cracks as shown in FIG. 22 are present in the nitrided material for 9 hours, and considering the corrosion of Mo at the tip of the crack, the corrosion rate of the Mo nitride layer is actually considered to be lower. . However, the SEM photographs before and after the corrosion of the nitride material in 50% nitric acid for 9 hours shown in FIG. 32 show that the Mo nitride layer was also completely corroded (Photo (b)). In order to improve the corrosion resistance of the Mo nitride layer against nitric acid, alloying with another passive metal is considered necessary.

Next, the corrosion resistance to hydrochloric acid solution was examined. FIG. 33 shows a comparison of corrosion rates of various metals in a hydrochloric acid solution. Ta, W, Mo, Zr are compared with other metals,
Excellent corrosion resistance. FIG. 34 shows a comparison of corrosion rates of an unnitrided Mo material and a nitrided material in NH ₃ gas at 950 ° C. for 9 hours in a 35% hydrochloric acid solution. This nitride material is described above (FIG. 23,
24, 25, 26), the β-Mo ₂ N phase is formed relatively thick on the sample surface. From FIG. 34, it can be seen that the non-nitrided Mo material has excellent corrosion resistance with a corrosion rate of 0.0059 mm / y, but the nitrided material has significantly improved corrosion resistance than this Mo material. That is, it was found that the corrosion rate of the nitrided material was 0.0011 mm / y, which was about 1/5 that of the unnitrided Mo material, and that the material had surprisingly high corrosion resistance.

Next, in order to examine the corrosion resistance of the Mo-based material in an alkaline solution, a corrosion test in a sodium hydroxide solution was performed. One example of the results is shown in FIG. FIG. 35 shows a comparison of corrosion rates of a non-nitrided Mo material and a material nitrided at 950 ° C. for 9 hours in NH ₃ gas in a 20% sodium hydroxide solution. Corrosion rate of unnitrided Mo material is 0.00
Although it shows complete corrosion resistance of 20 mm / y, the corrosion rate of the nitrided material is 0.0015 mm / y, indicating that the corrosion resistance is superior to that of the unnitrided Mo material.

(Experimental Example 3) Pure Mo nitrided as shown in FIG.
(A) and the optical microscope photograph of the cross section of Mo-1.0Ti alloy (b) are shown. After the cross section was filled with resin, mechanical polishing was performed, and then buff polishing was performed using a diamond slurry having a particle diameter of 1 μm. In each of the samples, the formation of a surface nitride layer having a thickness of about 80 μm was observed on the surface. Pure Mo
(FIG. 36 (a)) shows no difference in the polishing state of the structure inside the surface nitride layer, whereas the Mo-1.0T
In the i-alloy FIG. 36 (b), a clear difference is observed in the polishing state of the structure inside the surface nitride layer. In the region from the interface between the surface nitride layer and the mother phase to about 200 μm, the hardness is higher than that in the inside, so that the polishing rate is slow and the structure is flat with few irregularities. This region is an internal nitride layer in which it is considered that solid-solution Ti is preferentially nitrided and fine Ti nitride particles are dispersed and precipitated. In addition, Mo-1.0T
Even when the i-alloy was nitrided with NH ₃ gas, cracks were similarly observed in the surface nitride layer. FIG. 37 shows the nitrided pure Mo (a), Mo-0.5Ti (b) and Mo-
3 shows a hardness distribution of a cross section of a 1.0Ti (c) alloy. High hardness area (Hv-170) up to about 80 μm deep from the surface
0) corresponds to the Mo ₂ N surface nitride layer shown in FIG. When the Mo—Ti alloy is nitrided ((b),
(C), an internal nitride layer having high hardness, which cannot be seen in the case of pure Mo, is formed immediately inside the surface nitride layer. The hardness in this region is such that Hv is 800 to 9 immediately below the surface nitride layer.
It shows a high value of 80 and gradually decreases toward the inside. The maximum value of the hardness in the internal nitride layer is Hv-800 for the Mo-0.5Ti alloy and Hv-980 for the Mo-1.0Ti alloy, which depends on the amount of Ti. This high hardness in the internal nitride layer is presumed to be due to the dispersion strengthening of fine Ti nitride particles.

Regarding the strengthening of the Mo alloy by particle dispersion, a study in which TiC particles are dispersed by a mechanical alloying (MA) method is known. The maximum hardness of this alloy is between Hv and 500, and this value is based on TiC particles (1
wt%) and work hardening by MA. On the other hand, the strengthening of the internal nitride layer according to the present invention does not include work hardening and is purely due to the dispersion strengthening of Ti nitride particles. is there.

FIG. 38 shows the nitrided Mo-0.5Ti (a).
And the depth at which the hardness is maximum in the internal nitride layer of the Mo-1.0Ti (b) alloy (approximately 8
The incidence direction of the O-electron beam showing a transmission electron microscope photograph at 0 μm (about 130 μm for Mo-1.0Ti) is parallel to the [001] of the parent phase Mo in each case. It is clearly observed that a large number of particles having a strain field (corresponding to the black portions) (white elongated portions at the center of the strain field) are dispersed and precipitated.
As a result of electron beam diffraction, streaks were observed in the direction perpendicular to the particles extending in the <100> direction of the parent phase, these Ti nitride particles were deposited on the {100} plane of the parent phase Mo. It was found that the particles were precipitated as particles. The size is extremely fine and ultra-thin plate-like particles having a width of about 2 to 4 nm and a thickness of about 0.45 nm, and the particles are considered to be Ti nitride (TiN). Here, the thickness of the precipitate of about 0.45 nm corresponds to the size (a) of one unit cell of TiN.
_(TiN = 0.424 nm), and there is hardly any report on the precipitation and dispersion of such ultra-thin plate-like particles even with other metal materials. Also, compared to the Mo-0.5Ti alloy (a),
The Mo-1.0Ti alloy (b) having a larger amount has a larger number of precipitated particles. From these results, it can be seen that the hardening in the internal nitride layer is caused by the dispersed precipitation of ultrafine Ti nitride having a strain field. It is determined that there is.

FIG. 39 shows an HR-TEM image (a) of the precipitate dispersed in the internal nitride layer and its inverse Fourier transform (IF).
FT) Image (b) is shown. The sample surface was (00) as in FIG.
1) plane, the (110) plane and (1-10) of the parent phase
Plaid on the surface is visible. In the image (b) obtained by enlarging the particles indicated by ↑ in FIG. 39A and performing the inverse Fourier transform, considerable lattice distortion was observed around the precipitate. However,
These precipitates are connected to the matrix and the lattice, and maintain perfect consistency. No dislocation due to misfit between the matrix and the lattice of the precipitate is observed. Since the solute element is Ti, the precipitate is considered to be TiN having an fcc structure. The TiN particles have a size of $ 10 of the parent phase Mo (bcc).
When deposited on the 0 ° plane, each lattice constant is a _Mo
= 0.314 [nm], a _TiN = 0.424 [nm]
Therefore, the azimuth relationship that minimizes the misfit is the relationship of the following equation as shown in FIG.

[0057]

(Equation 4)

[0058]

(Equation 5) This relationship is based on the backer notting (Backer-notching) often seen in the case of a precipitate in a γ-Fe alloy (bcc).
(Nuting). This orientation relationship is also consistent with the result of internal nitriding of the Mo-Ti alloy at 1300 ° C in pure nitrogen.

When TiN precipitates while maintaining perfect coherence with the parent phase, the maximum size of its diameter D and thickness t is equal to the gaps δD and δt of the misfit dislocations. )it is conceivable that.

[0060]

(Equation 6)

[0061]

(Equation 7) Where ε ₁ = | d _{(220) Mo} −d _{(200) TiN}
|, Ε ₂ = | d _{(200) Mo−} d _{(200) TiN} |
It is. Therefore, d _{(200) Mo} = 0.15n
m, d _{(220) Mo} = 0.222 nm, d
_When the calculation is performed by substituting _{(200) TiN} = 0.212 nm, the size of TiN that can be precipitated while maintaining the consistency is:

[0062]

(Equation 8)

[0063]

(Equation 9) It can be estimated. As described above, the precipitate size observed in the present invention has a width of about 2 to 4 nm and a thickness of 0.45 nm, and this observed value almost coincides with the calculated value of the formula (6.7). Therefore, in the internal nitride layer when nitrided at 1100 ° C., a unit (0.424 nm) of Ti
It is considered that N was precipitated while maintaining consistency.

The fine distribution of precipitates at various depths in the inner nitrided layer of the nitrided Mo-1.0Ti alloy was determined by TEM.
Was observed directly. In each case, the incident direction of the electron beam is parallel to the parent phase Mo [001] as in FIG. FIG.
(A), (b) and (c) show that the depth d from the sample surface of the Mo-1.0Ti alloy is d = 130 μm, the position of the maximum value (値 980), and the intermediate value (〜700). It corresponds to the reduced position and the position reduced to a value close to the value of the parent phase (（310). From these photographs, it was clarified that the density of the TiN precipitate particles decreased as the position was more inside the sample, that is, as the hardness decreased. It was also found that the size of the precipitate particles hardly changed. Therefore, it can be considered that a considerable amount of solid solution Ti still exists in the low hardness regions in FIGS. 41 (b) and 41 (c).

T measured from a photograph of the TEM observation in FIG.
FIG. 42 shows the relationship between the square root (N) of the distribution density of iN precipitate particles and the amount of increase in hardness (△ Hv). Here, △ Hv is the hardness at each depth Hv The hardness of the parent phase Mo (Hv)
The difference of the _matrix . That is,

[0066]

(Equation 10) From this it can be seen that ΔHv is proportional to the square root of N. That is,

[0067]

[Equation 11] Here, when the shape of the precipitate is approximated by a cube as a first approximation, the volume fraction (f) is the volume (V) of the matrix, the particle size (r), and the number of particles (n) contained therein. ) Can be expressed by the following equation.

[0068]

(Equation 12) Considering the result that no change was observed in the size of the precipitate in the internal nitride layer in the TEM observation in FIG. 41, it can be considered that the size r of the precipitate is constant in the above equation. In this case, the following relationship is obtained between f and N.

[0069]

(Equation 13) From Equations (13) and (11), it can be seen that ΔHv is proportional to the square root of the volume fraction f of the TiN precipitate particles. That is,

[0070]

[Equation 14] From the above examination, it has been clarified that the decrease in ΔHv in the depth direction in the internal nitride layer can be understood by the decrease in the volume fraction of TiN particles in the depth direction as shown in Expression (14). .

The mechanism of strengthening when the second phase is precipitated by nitriding can be roughly classified into two types. That is, the precipitate is small, the dislocation can pass through the precipitate, the shear mechanism of the precipitate by the dislocation, and the dislocation loop leaving around the precipitate when the precipitate is large and the dislocation bypasses the precipitate. There are two types of the Orowan mechanism that will resist the movement of the next dislocation. It has been experimentally known that ΔHv is proportional to the square root of the volume fraction of precipitates in the case of a shear mechanism by dislocation of fine precipitate particles having matching strain. The precipitate in the present invention is a very small and thin particle that maintains consistency with the parent phase, and it is considered that the dislocation can be sufficiently passed.
Considering the proportional relationship between ^1/2 , the mechanism of strengthening the internal nitride layer by nitriding at 1100 ° C. is determined to be a mechanism of shearing precipitates by dislocation.

As described above, when the Mo—Ti alloy is nitrided by using NH ₃ gas, the inside of the surface Mo nitride layer has a high phase hardness, which is presumed to be due to the dispersion and precipitation of fine TiN particles. A nitride layer is formed. The hardness in the internal nitride layer has a maximum value immediately below the surface nitride phase, and gradually decreases as it goes inward. The maximum value of the hardness of the internal nitride layer is Hv-800, Mo-1.
Hv-980 for the 0Ti alloy, and the higher the Ti content, the harder.

[0073]

[Example] Mo-0.5wt% Ti prepared by powder metallurgy
Rolled alloy material and cut square bar-shaped sample (2 mm x 1
mm × 25 mm) in a stream of NH ₃ gas of 1 atm.
After nitriding at 0 ° C. for 0.5 hour, β-Mo _{2 having} a thickness of 7 μm
An N layer was obtained. As a corrosion resistance test, the processed product was immersed in dilute sulfuric acid for one week and then taken out and observed, but no corrosion was found. The hardness in this case was Hv-800. Sho Mo
Has a hardness of Hv-200.

[0074]

According to the present invention, Mo or a Mo-based alloy is
° C. and treated for 0.2 hour to 100 hours in an atmosphere of to 1150 ° C., since the base material surface provided Mo ₂ N layer 0.5 ~ 10 m, with dramatically improve the corrosion resistance, significantly improved the hardness, Compared with Ta in terms of corrosion resistance performance, it has various effects such as higher mechanical strength than Ta, lighter weight than Ta, and lower price.

[Brief description of the drawings]

FIG. 1 shows pure Mo and Mo-0.5Ti alloy at 1100 ° C.
The figure which shows the XRD pattern of the surface of the sample nitrided for 16 hours at FIG.

FIG. 2 shows a method of (a) 600 ° C. to (f) 11 using a Mo—Ti alloy.
At 00 ° C., shows an XRD pattern in the 16 hours N ₂ gas nitriding treated sample surface.

FIG. 3 (a) Mo-0.5Ti alloy was heated at 1100 ° C. for 4 hours.
Photomicrograph of metal structure of cross-section after time nitriding. (B) An optical microscope photograph of the metal structure of the cross section also subjected to nitriding treatment for 16 hours.

FIG. 4 is a graph showing the relationship between the thickness of the surface nitride layer of the Mo-0.5Ti alloy and the nitriding time.

FIG. 5 is a graph showing a state in which the thickness of a sample does not change during nitriding of Mo-0.5Ti.

FIG. 6 is a graph showing surface and internal hardness distributions when a sample of pure Mo (a) and Mo-0.5Ti (b) is nitrided at 1100 ° C. for 16 hours.

FIG. 7 shows a depth d = 0 when a sample of pure Mo, β-Mo ₂ N and γ-Mo ₂ N is subjected to a nitriding treatment at 1100 ° C. for 16 hours.
(A)-The graph which shows the crystal structure in d = 80 (f).

FIG. 8 is a graph showing a change in lattice constant of Mo nitride with depth obtained by X-ray diffraction.

FIG. 9 is a graph showing the axial ratio (c / a) of Mo nitride obtained from the result of X-ray diffraction of pure Mo nitrided for 16 hours.

10A is a schematic diagram of a crystal structure of γ-Mo ₂ N. FIG. (B) Schematic diagram of the crystal structure of β-Mo ₂ N.

FIG. 11 is a transmission electron microscope photograph of a β-phase metal surface formed inside a surface nitride layer (depth: 60 μm) of a sample obtained by nitriding a Mo-0.5Ti alloy for 16 hours.

FIG. 12 is a transmission electron micrograph of a metal surface of β-Mo ₂ N subjected to electron diffraction when Mo was nitrided at 1100 ° C. for 16 hours.

13 (a) is an analysis diagram of [100] incidence of the sample of FIG. (B) Analysis diagram of [211] incidence of the sample in FIG.

FIG. 14 is a diagram showing that there are two types of lamination methods of β phase.

FIG. 15A is a diagram schematically showing an atomic arrangement. (B) A diagram showing that the atomic arrangement is mirror symmetric.

FIG. 16 is a high-resolution electron microscope photograph of a β-phase metal structure photographed at a magnification of 550,000 times.

FIG. 17 is an inverse Fourier transform image obtained by enlarging FIG. 16 and shows a high-resolution electron microscope photograph of a metal structure.

FIG. 18 (a) is an inverse Fourier transform image of a stretched image of FIG. 16, which is a high resolution electron microscope image of a metal structure. (B) The same figure as (b) of FIG.

FIG. 19 is a conceptual diagram of a corrosion test device.

FIG. 20 is a diagram showing the degree of corrosion of various metals.

FIG. 21 shows a Mo-0.5Ti alloy in an NH ₃ gas at 95%.
It is a SEM photograph of the metal structure of the surface before and after the corrosion test in boiling sulfuric acid when nitriding treatment at 0 ° C. for 9 hours,
(A) is a sulfuric acid concentration of 40%, (b) is 60%, and (c) is 7%.
5%.

FIG. 22 is a graph showing the corrosion rate of a sample nitrided at 950 ° C. for 9 hours in boiling sulfuric acid.

FIG. 23 (a) is an optical microscope photograph of a metal structure on a sample surface before corrosion of a sample which was nitrided at 950 ° C. for 9 hours using a 10 mm square sample. (B) An optical microscope photograph of the metal structure on the surface of the strip-shaped sample recrystallized before nitriding.

FIG. 24 shows a Mo-0.5Ti alloy in NH ₃ gas at 95%.
75% before corrosion (a) of sample nitridated at 0 ° C for 9 hours
Optical micrograph of the metallographic structure after (b) corrosion test in concentrated sulfuric acid for 8 days.

FIG. 25 (a) is an optical microscope photograph of the metal structure on the surface of a sample obtained by nitriding the recrystallized material at 950 ° C. for 1 hour. (B) An optical microscope photograph of the metal structure on the surface of the sample obtained by nitriding the recrystallized material at 950 ° C. for 30 minutes.

FIG. 26A shows a nitrided material for 9 hours. (B) 1 hour nitriding material. (C) a nitride material for 30 minutes. 4 is a graph showing an XRD pattern of the above surface.

FIG. 27 is a graph showing the corrosion rates of pure Mo when nitrided and unnitrided.

FIG. 28 is an optical microscope photograph of a metal structure on a sample surface of a nitride material for 30 minutes in a corrosion test performed in 75% boiling sulfuric acid for 5 days.

FIG. 29 is a view showing an XRD pattern on the surface of a sample obtained by nitriding pure Mo in a N ₂ gas at 800 ° C. for 48 hours.

FIG. 30 is a graph showing corrosion rates of various refractory metals in boiling nitric acid.

FIG. 31 shows pure Mo (a) and M subjected to nitriding treatment for 9 hours.
Graph of a corrosion test for nitric acid with o ₂ N (b).

FIG. 32: SE of metal surface before corrosion (a) and before corrosion (b) in 50% nitric acid of Mo ₂ N (9 hours nitride)
M Shashin.

FIG. 33 is a graph showing the corrosion rate in 75% boiling sulfuric acid.

FIG. 34 is a graph showing the corrosion rate in 35% boiling sulfuric acid.

FIG. 35 is a graph showing the corrosion rate in 20% sodium hydroxide.

FIG. 36 is an optical microscope photograph of the metal structures of pure Mo (a) and Mo-0.5Ti alloy (b).

FIG. 37 is a diagram showing the hardness distribution of nitrided pure Mo (a), Mo-0.5Ti alloy (b), and Mo-1.0Ti alloy (c).

FIG. 38 shows a nitrided pure Mo-0.5Ti alloy (a);
A transmission electron microscope photograph of the metal structure at a depth at which the hardness of the Mo-1.0Ti alloy (b) is maximized.

FIG. 39 is a transmission electron microscope photograph of a metal structure showing an HR-TEM image (a) and an inverse Fourier transform (IFFT) image (b) of a precipitate dispersed in a nitride layer.

FIG. 40 is a view showing an azimuth relationship in which misfit is minimized from a lattice constant when TiN particles are precipitated on a mother phase Mo surface.

FIG. 41 shows that the depth d from the sample surface of the Mo-1.0 Ti alloy is d = 130 μm (a) and d = 225 μm, respectively.
(B), Photograph of metal structure observed by TEM at d = 300 μm (c).

42 is a graph showing the relationship between the distribution density N of TiN precipitate particles measured from the photograph of FIG. 41 and the increase in hardness (ΔHv).

[Table 5]

Claims (5)

[Claims] The thickness of the Mo alloy is 0.5 μm to 10 μm.
High corrosion resistance M characterized by providing a Mo ₂ N layer of μm
o-based composite materials. 2. The high corrosion resistant Mo-based composite material according to claim 1, wherein the Mo ₂ N layer is a β-Mo ₂ N layer. 3. A highly corrosion-resistant Mo-based composite characterized by subjecting a Mo-based alloy to nitriding treatment in an atmosphere of 700 ° C. to 1150 ° C. for 0.2 to 100 hours in the presence of N ₂ gas or NH ₃ gas. Material manufacturing method. 4. A Mo-based alloy is nitrided in an atmosphere of 700 ° C. to 1150 ° C. in the presence of N ₂ gas or NH ₃ gas, and has a thickness of 0.5 μm or more and which does not cause cracks.
A method for producing a highly corrosion-resistant Mo-based composite material, comprising providing an o ₂ N layer. 5. The method for producing a highly corrosion-resistant Mo-based composite material according to claim 4, wherein the thickness not causing cracks is less than 10 μm.