JP5157076B2 - Method for producing sintered body of magnetic alloy - Google Patents

Description

  The present invention relates to a magnetic alloy material suitably used as a magnetic refrigeration working substance or a magnetostrictive material and a method for producing the same.

In recent years, composition formula: La _1-z RE _z (Fe _1-x A _xy TM _y ) ₁₃ (A = at least one element of Al, Si, Ga, Ge, Sn 0.05 ≦ x ≦ 0.2) , TM is a magnetic alloy represented by at least one element of transition metal elements 0 ≦ y ≦ 0.1, RE is at least one element of rare earth elements excluding La, 0 ≦ z ≦ 0.1) Hereinafter, “La (Fe, Si) ₁₃ -based magnetic alloy” has a NaZn ₁₃ type crystal structure, and exhibits a large magnetocaloric effect and magnetovolume effect near the Curie temperature (Tc). It is considered promising as a magnetic refrigeration working substance and a magnetostrictive material (for example, Patent Document 1, Patent Document 2).

Conventionally, a La (Fe, Si) ₁₃ series magnetic alloy is produced by, for example, heat-treating a die casting alloy obtained by arc melting or high frequency melting in vacuum at 1050 ° C. for about 168 hours, The production efficiency was very low.

The present applicant disclosed in Patent Document 3 a method for producing a La (Fe, Si) ₁₃ -based magnetic alloy material with high efficiency by using a liquid quenching method (also referred to as a superquenching method). However, when the magnetic alloy material disclosed in Patent Document 3 is used as a magnetic refrigeration working material, a magnetic material having a large thermal contact area with a heat exchange medium using an alloy material obtained by roughly pulverizing a ribbon-like alloy material. It is necessary to prepare a frozen working substance. For example, as the heat exchange medium, a liquid heat exchange medium such as a water-based antifreeze liquid having a relatively high specific heat and good fluidity at an operating temperature or a low-melting hydrocarbon solvent is preferably used as a magnetic refrigeration working material. May be a bed in which a coarse powder of an alloy material is stored in a bowl-shaped container, a compact formed by compression molding into a thin plate, or a sintered body sintered in a porous bulk shape through which liquid can pass. .

On the other hand, Patent Document 4 describes a method for obtaining a La (Fe, Si) ₁₃ -based magnetic alloy sintered body having a desired shape by using powder metallurgy. In powder metallurgy, a sintered compact is obtained by sintering a compact (compact) obtained by press-molding an alloy powder (fine powder). The degree is large, and the processing cost can be reduced instead.
JP 2000-54086 A JP 2002-69596 A JP 2004-100043 A JP 2005-36302 A

As a result of studying the application of the powder metallurgy method to the La (Fe, Si) ₁₃ -based magnetic alloy material, the present inventor has found the following problems.

In order to apply the powder metallurgy method to the La (Fe, Si) ₁₃ -based magnetic alloy material disclosed in Patent Document 3, a NaZn ₁₃ type compound phase as a target phase is generated by heat treatment, and then finely pulverized. It is necessary to produce a molded body using the obtained powder and to sinter it. That is, although the heat treatment time can be greatly shortened by using the manufacturing method described in Patent Document 3, a heat treatment for obtaining a NaZn ₁₃ type compound phase and a heat treatment for sintering are performed twice in total. Must be carried out in a vacuum, and the production efficiency is poor.

In addition, the method described in Patent Document 3 generates a NaZn ₁₃ type compound phase by a solid phase reaction based on an element diffusion process in a ribbon of an as-spun alloy (quenched alloy), and therefore includes a relatively large structure. However, a NaZn ₁₃ type compound phase can be obtained by heat-treating the quenched alloy ribbon. However, when the powder metallurgy method is applied to a quenched alloy having a large structure, each phase constituting the structure becomes a separate particle, or the composition varies greatly among the particles. This requires elemental movement between powder particles, which requires long-time sintering (heat treatment), which is not practically preferable.

  Furthermore, the finely pulverizing property in the as-spun state is generally very poor in a structure in which coarse primary crystal iron has developed into a columnar shape, and even when the ultra-quenching method is applied, the particle size required for powder metallurgy (for example, 2 μm). ), If the iron primary crystal phase is larger than that, it is extremely difficult to obtain a powder having a target particle size.

The alloy raw material described in Patent Document 4 is also produced under a slow cooling rate of 1 × 10 ⁴ ° C./s, so that the structure is not sufficiently fine. As a result, in order to obtain a sintered body substantially composed of the NaZn ₁₃ type compound phase, (1) the ratio of the NaZn ₁₃ type compound phase in the alloy needs to be 85 mass% or more in advance by heat treatment of the raw material alloy. There are problems such as (2) requiring a long sintering time and (3) requiring a high sintering temperature of 1280 ° C. or higher.

  A problem in the sintering process when the mother alloy is an alloy with a low quenching degree such as an ingot is that the α-Fe phase cannot be effectively lost at a sintering temperature lower than the peritectic temperature. On the other hand, above the peritectic temperature, a new α-Fe phase or LaFeSi compound phase is generated. Therefore, in order to obtain a single-phase and high-density sintered body, it has been necessary to perform sintering for a long time at a high temperature and in a narrow temperature range.

The present invention has been made in view of the above points, and the main object of the present invention is to sinter a sintered body containing a NaZn ₁₃ type compound phase by using powder metallurgy method at a relatively low cost and in a short time. It is in providing the raw material alloy (powder) used for the manufacturing method and its manufacturing method.

Magnetic alloy material of the present invention, the composition formula _{Fe 100-abc RE a A b} Co c (RE is a rare earth element necessarily including La, A is Si or Al, 6at% ≦ a ≦ 11at %, 8at% ≦ b ≦ 18at %, 0 at% ≦ c ≦ 9 at%), and is substantially a two-phase structure consisting of an α-Fe phase and a (RE, Fe, A) phase containing 30 to 90 at% RE, or α A three-phase structure comprising an Fe phase, a (RE, Fe, A) phase containing RE of 30 at% to 90 at% and a RE (Fe, A) ₁₃ compound phase of NaZn ₁₃ type crystal structure, The average size in the minor axis direction is in the range of 40 nm to 2 μm. The magnetic alloy material of the present invention is an as-spun alloy (an alloy that has not undergone heat treatment after rapid solidification) produced by a liquid quenching method (including a strip casting method and a melt spinning method).

  In one embodiment, a in the composition formula is 7 at% or more and 9 at% or less.

  In one embodiment, the (RE, Fe, A) phase is a REFeSi compound phase.

In one embodiment, the Co substitutes Fe in at least one of the α-Fe phase, the (RE, Fe, A) phase, and the RE (Fe, A) ₁₃ compound phase.

  In one embodiment, the oxygen content is 0.07 atomic% or more and 0.18 atomic% or less.

In some embodiments, substantially all can be converted to the RE (Fe, A) ₁₃ compound phase by heat treatment at a temperature of 600 ° C. or higher for a time of 10 seconds or longer.

  The magnetic alloy material of an embodiment is a powder, and the size of the minor axis direction of the particles constituting the powder is in the range of 2 μm or more and 200 μm or less.

  In one embodiment of the magnetic alloy material, the size of the particles constituting the powder in the minor axis direction is less than 10 μm.

The method for producing a magnetic alloy material according to the present invention includes a step of preparing a molten alloy raw material having a predetermined composition, and an average cooling rate in the temperature range of 1500 ° C. to 600 ° C. of the molten alloy raw material is 2 × 10. at ⁴ ° C. / s to 2 × a 10 ⁶ ° C. / s within the range of conditions by rapid solidification, the composition formula _{Fe 100 - abc RE a a b} Co c (RE is a rare earth element necessarily including La, A is represented by Si or Al, 6 at% ≦ a ≦ 11 at%, 8 at% ≦ b ≦ 18 at%, 0 at% ≦ c ≦ 9 at%), and is substantially an α-Fe phase and a RE of 30 at% to 90 at%. A two-phase structure composed of (RE, Fe, A) phase containing Na or a substantially α-Fe phase, (RE, Fe, A) phase containing 30 at% to 90 at% RE and NaZn ₁₃ type crystal RE (Fe, a) ₁₃ of the structure Has a 3-phase structure consisting of phase, the average size of each phase in the short axis direction is characterized in that it comprises a step of forming the rapidly solidified alloy is in the range of 40nm or more 2μm or less.

  In one embodiment, a in the composition formula is 7 at% or more and 9 at% or less.

  In one embodiment, the (RE, Fe, A) phase is a REFeSi compound phase.

In one embodiment, the Co substitutes Fe in at least one of the α-Fe phase, the (RE, Fe, A) phase, and the RE (Fe, A) ₁₃ compound phase.

  In one embodiment, the thickness of the quenched alloy is 2 μm or more and 200 μm or less.

  In one embodiment, in the step of forming the quenched alloy, the tapping temperature of the alloy raw material is higher than the liquidus temperature of the alloy raw material by a temperature of 50 ° C. or higher and 150 ° C. or lower.

  In one embodiment, in the step of forming the quenched alloy, a roll peripheral speed of the cooling roll is in a range of 3 m / s to 30 m / s.

  In one embodiment, the method further includes a step of pulverizing the quenched alloy to obtain a powder composed of particles having a minor axis size in the range of 2 μm to 200 μm.

  In a certain embodiment, the size of the minor axis direction of the particles constituting the powder is less than 10 μm.

  The method for producing a sintered body of a magnetic alloy according to the present invention includes a step of preparing a powder obtained by the above production method, a step of obtaining a molded body by molding the powder, and sintering the molded body. And a process.

  In one embodiment, the sintering step includes a step of sintering within a temperature range of 600 ° C. or higher and lower than 1320 ° C. In some embodiments, the sintering temperature is 900 ° C. or higher.

  In one embodiment, in the sintering step, the time for sintering within the temperature range is not less than 10 seconds and not more than 8 hours. A more preferable sintering time is 4 hours or less.

Since the average size in the minor axis direction of each phase constituting the La (Fe, Si) -based magnetic alloy of the present invention is in the range of 40 nm or more and 2 μm or less, the inter-particles in the sintering process when using powder metallurgy method Thus, the La (Fe, Si) ₁₃ -based NaZn ₁₃ type compound phase, which is the target phase, can be obtained in a short time. That is, according to the present invention, a La (Fe, Si) ₁₃ -based NaZn ₁₃ type compound phase can be obtained by a single heat treatment (sintering) using powder metallurgy.

In addition, the magnetic alloy of the present invention can be easily pulverized in an as-spun (quenched alloy) state, and the raw material powder obtained by pulverization has a relatively low oxygen concentration, and is produced using the raw material powder. The oxygen concentration of the (La, Fe, Si) -based magnetic alloy (sintered body) obtained by sintering (heat-treating) the formed body is relatively low. Therefore, the magnetic properties of the La (Fe, Si) ₁₃ type compound obtained by pulverizing, forming and sintering the as-spun alloy are all equal to or higher than those of the conventional, and as a magnetic refrigeration working material or magnetostrictive material. Used effectively.

The inventors of the present invention applied the powder metallurgy method to the quenched alloy disclosed in Patent Document 3 by using a quenched alloy (as-spun alloy) before heat treatment as a raw material alloy, thereby forming a final shape. In this case, by controlling the microstructure of the raw material alloy sufficiently, the elements inside the powder particles do not require element diffusion between the powder particles having a long time. It was thought that the NaZn ₁₃ type compound phase, which is the target phase, could be obtained by the sintering process only by diffusion.

Therefore, as a result of intensive studies, the present inventor has found an as-spun alloy material having a structure suitable for a raw material powder used for producing a molded body of a La (Fe, Si) ₁₃ -based magnetic refrigeration working material.

[Basic composition]
Alloy magnetic material according to the present invention, the composition formula _{Fe 100 - abc RE a A b} Co c (RE is a rare earth element necessarily including La, A is Si or Al, 6at% ≦ a ≦ 11at %, 8at% ≦ b ≦ 18 at%, 0 at% ≦ c ≦ 9 at%), and substantially a two-phase structure consisting of an α-Fe phase and a (RE, Fe, A) phase containing 30 at% to 90 at% RE, or substantially Has a three-phase structure consisting of an α-Fe phase, a (RE, Fe, A) phase containing RE of 30 at% to 90 at% and a RE (Fe, A) ₁₃ compound phase of NaZn ₁₃ type crystal structure, The magnetic alloy material has an average size in the minor axis direction of each phase in the range of 40 nm or more and 2 μm or less. RE is selected from the group consisting of La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Dy, Ho, Er, and Tm and is at least one rare earth element containing 90 at% or more of La Is preferred. A preferably contains Si, and more preferably Si.

The La (Fe, Si) ₁₃ type compound has a wide ratio of Fe to Si as shown in the ternary phase diagram of La—Fe—Si at 600 ° C. in FIG. 24, but is suitable for a magnetic refrigeration working material. The ratio is the composition on the Fe-rich side. An object of the present invention is to finally obtain a sintered body which is a NaZn ₁₃ type compound and has a compound composition of Fe-rich.

Taking the ternary composition of La-Fe-Si as an example, among several La-Fe-Si compounds, it is desirable to have a LaFeSi compound containing La of 30 at% or more in order to obtain the target compound. On the other hand, when a LaFe ₂ Si ₂ compound containing La less than 30 at% is produced, the ratio of Fe to Si in the obtained NaZn ₁₃ type compound is shifted to the Si-rich side, and the magnetic refrigeration working material is Performance decreases. Therefore, the RE of the (RE, Fe, A) phase that should coexist with α-Fe needs to be 30 at% or more.

On the other hand, depending on the alloy composition and the rapid cooling process, a eutectic structure composed of a La ₅ (Fe, Si) ₃ compound phase, metal La, and metal Fe may be formed as the (RE, Fe, A) phase. The crystal structure contributes to the formation of La (Fe, Si) ₁₃ type compound as the target phase by generating LaFeSi in the heat treatment process. This eutectic structure is recognized as a rare earth-rich region in a gap in which an α- (Fe, Si) phase, which is a primary crystal composed of a plurality of phases, has grown, and depending on the analysis technique, one rare earth-rich phase (RE, Fe, It may be recognized as A). The average composition varies depending on the cooling conditions during solidification, but the La concentration does not exceed about 90 at%, which is a Fe—La eutectic composition. Therefore, the upper limit of the RE concentration of the (RE, Fe, A) phase that should coexist with α-Fe is 90 at%.

Also, La (Fe, Si) ₁₃ is not generated unless α- (Fe, Si), LaFeSi, and La (Fe, Si) ₁₃ are in the composition region. In order to satisfy this condition, it is necessary to set a lower limit for the Si concentration. Below this lower limit, α- (Fe, Si), La ₅ (Fe, Si) _3, and LaFeSi enter a composition region in which they coexist and La (Fe, Si) ₁₃ is not generated. Therefore, since the lower limit value of RE is 6%, the lower limit value of A is about 8 at%. As will be described later in detail in the examples, the lower limit value of RE is preferably 7%, and the lower limit value of A at that time is about 8.3%. Furthermore, when the Si concentration increases, the Curie temperature of the La (Fe, Si) ₁₃ phase finally generated increases, but at the same time, the saturation magnetization and the magnetocaloric effect are reduced. When the Si concentration exceeds 18 at% (x = 0.167), the value of −ΔSmag between the applied magnetic field 0T and 1T of the alloy is less than 1 J / K / kg, and the practical value is lost. %.

  The composition ratio of RE will be described in detail later in Examples.

  The magnetic alloy material (raw material alloy) according to the present invention is excellent in fine pulverization in an as-spun state since the solidification primary crystal mainly composed of a soft iron phase that inhibits pulverization because it is viscous and rich in toughness is 2 μm or less. Since the element diffusion distance is short in the sintering reaction after molding, a sintered body having excellent composition uniformity can be obtained by short-time sintering. As a result, it is possible to obtain a magnetic refrigeration working material that is easy to give form as a magnetic refrigeration working material, has a sharp magnetization temperature change at the magnetic transition temperature, and has a large magnetocaloric effect. Of course, the sintered body obtained by the powder metallurgy process using the magnetic alloy material (raw material alloy) according to the present invention is also suitably used as a magnetostrictive material.

[Size of crystal grains in the minor axis direction]
The upper limit of the crystal grain size is determined by the diffusion distance during sintering (heat treatment). For example, in order to generate a target phase by heat treatment at 900 ° C. for about 1 hour, the size of the structure must be about 3 μm or less. Further, in order to ensure the uniformity of the composition of the product phase, it must be finer and further smaller than the powder particle size required in the powder metallurgy process. The upper limit of the crystal grain size that satisfies these requirements is 2 μm. That is, when the size of the produced structure exceeds 2 μm, the efficiency of the pulverization process is extremely reduced, and the sintering time for obtaining the target phase or the heat treatment time after pulverization and powder press molding becomes too long.

  On the other hand, the lower limit crystal grain size that can be practically produced is 40 nm to 50 nm. If crystal grains having a particle size smaller than this are made, the quenched alloy becomes amorphous. When it becomes amorphous, it is easy to obtain a thin or powdery material, but it is difficult to grind it in an amorphous state and it is not suitable for a powder metallurgy process.

  The crystal grain size is almost determined by the cooling rate. The larger the cooling rate, the smaller the crystal grain size and the thinner the alloy. The presence of a lower limit on the thickness of the quenched alloy results in an upper limit on the cooling rate, and the secondary result determines the lower limit of the crystal grain size.

  As will be shown later, as a result of various experiments, by optimizing the conditions of the liquid quenching method, an alloy having an average size in the minor axis direction of each phase constituting the quenched alloy in the range of 40 nm to 2 μm. Can be obtained.

[Oxygen concentration]
Further, the oxygen concentration of the magnetic alloy material (raw material alloy before sintering) of a preferred embodiment according to the present invention is 0.02 mass% or more and 0.05 mass% or less (0.07 atom% or more and 0.18 atom% or less). In addition, the oxygen concentration of a sintered body (magnetic refrigeration material) produced using such a magnetic alloy material is lower than that of the conventional one, and as a result, the value of change in magnetic entropy is improved.

  A lower oxygen concentration is preferred. Oxygen is incorporated as an impurity in the raw material melting step, the fine pulverization step, and the sintering step, and mainly combines with La and partly combines with Si. The target substance changes its magnetic transition behavior depending on the quantity ratio between La and (Fe + Si) and the quantity ratio between Fe and Si, and the magnetocaloric effect is sensitively affected. Oxygen is inevitably taken up, but its concentration is difficult to control, and the higher the oxygen concentration, the greater the variation in magnetic properties. This results in a distribution of magnetic transition points within the material, which impairs the sharp transition and reduces the magnetocaloric effect.

  In order to obtain a good magnetocaloric effect, it is desirable that the oxygen concentration in the sintered body is 0.08 mass% or less, and in the raw material alloy, 0.05 mass% or less is preferable. Assuming that the oxygen concentration inevitably increases in the powder metallurgy process, the oxygen concentration of the magnetic alloy material (raw material alloy) of the present invention is 0.02 mass% or more and 0.05 mass% or less (0.07 atom). % Or more and 0.18 atomic% or less).

[Liquid cooling method (super rapid cooling method)]
The magnetic alloy material having a fine structure described above has an average cooling rate of 2 × 10 ⁴ ° C./s or more and 2 × 10 ⁶ ° C. in a temperature range from 1500 ° C. to 600 ° C. of a molten alloy material having a predetermined composition. It is produced by rapid solidification under conditions within the range of / s or less. The average cooling rate is preferably 2 × 10 ⁵ ° C./s or more.

As the liquid quenching method, a gas atomizing method, a single roll quenching method, a twin roll quenching method, a strip casting method, a melt spinning method, or the like can be used. In particular, use of a melt spinning method, a strip casting method, or the like is preferable because a ribbon-like quenched alloy having a thickness of 20 μm to 200 μm can be manufactured with high efficiency. The gas atomization method generally has a slow cooling rate and it is difficult to obtain an average cooling rate of 2 × 10 ⁴ ° C./s or more. However, if the particle size of the powder is small, for example, 50 μm or less, the degree of supercooling ( ΔT) becomes 100K or more, and a cooling rate of the order of 10 ⁵ ° C / s can be obtained, and spherical particles can be obtained. The powder composed of spherical particles is suitable as, for example, a raw material powder for a compound for injection molding because it is easy to mix and knead with a polymer binder and is excellent in moldability (fluidity) of the obtained compound.

  For example, a quenching alloy can be produced by a melt spinning method using the quenching apparatus shown in FIG. In order to prevent oxidation of a raw material alloy containing rare earth elements (La and RE in the above composition formula) and Fe that are easily oxidized, it is preferable to perform the quenching step in an inert gas atmosphere. As the inert gas, a rare gas such as helium or argon or nitrogen can be used. Note that since nitrogen is relatively easy to react with rare earth elements, it is preferable to use a rare gas such as helium or argon.

  The rapid cooling apparatus shown in FIG. 1 includes a raw material alloy melting chamber 1 and a rapid cooling chamber 2 capable of maintaining a vacuum or an inert gas atmosphere and adjusting the pressure. FIG. 1A is an overall configuration diagram, and FIG. 1B is a partially enlarged view.

  As shown in FIG. 1 (a), the melting chamber 1 includes a melting furnace 3 for melting a raw material 20 blended so as to have a desired alloy composition at a high temperature, and a hot water storage container 4 having a hot water discharge nozzle 5 at the bottom. And a blended raw material supply device 8 for feeding the blended raw material into the melting furnace 3 while suppressing the entry of air. The hot water storage container 4 has a heating device (not shown) that stores the molten alloy 21 of the raw material alloy and can maintain the temperature of the hot water at a predetermined level.

  The quenching chamber 2 includes a rotary cooling roll 7 for rapidly cooling and solidifying the molten metal 21 discharged from the hot water nozzle 5.

  In this apparatus, the atmosphere in the melting chamber 1 and the quenching chamber 2 and the pressure thereof are controlled within a predetermined range. Therefore, atmospheric gas supply ports 1b, 2b, and 8b and gas exhaust ports 1a, 2a, and 8a are provided at appropriate locations in the apparatus. In particular, the gas exhaust port 2a is connected to a pump in order to control the absolute pressure in the quenching chamber 2 within a range of 30 kPa to normal pressure (atmospheric pressure) (preferably 100 kPa or less). By changing the pressure in the melting chamber 1, the injection pressure of the molten metal coming out of the nozzle 5 can be adjusted.

  The melting furnace 3 can be tilted, and the molten metal 21 is appropriately poured into the hot water storage container 4 through the funnel 6. The molten metal 21 is heated in the hot water storage container 4 by a heating device (not shown).

  The hot water discharge nozzle 5 of the hot water storage container 4 is disposed in the partition wall between the melting chamber 1 and the quenching chamber 2 and causes the molten metal 21 in the hot water storage container 4 to flow down to the surface of the cooling roll 7 positioned below. The orifice diameter of the hot water nozzle 5 is set within a range of 0.5 mm to 4.0 mm. By adjusting the orifice diameter and / or the pressure difference between the melting chamber 1 and the quenching chamber 2 (for example, 10 kPa or more) according to the viscosity of the molten metal 21, the molten metal 21 is smoothly discharged. According to the apparatus used in the present embodiment, the supply rate of the molten alloy can be set to 1.5 kg / min to 10 kg / min. When the supply rate exceeds 10 kg / min, the molten metal cooling rate becomes slow, and there arises a disadvantage that a multilayer structure in which the structure changes in the thickness direction of the slab is formed. A more preferable supply rate of the molten alloy is 2 kg / min to 8 kg / min.

  The cooling roll 7 is preferably formed from Cu, Fe, or an alloy containing Cu or Fe. If the cooling roll is made of a material other than the alloy containing Cu or Fe, the peelability of the rapidly cooled alloy with respect to the cooling roll is deteriorated, so that the rapidly cooled alloy may be wound around the roll. The diameter of the cooling roll 7 is, for example, 300 mm to 500 mm. The water cooling capacity of the water cooling device provided in the cooling roll 7 is calculated and adjusted according to the solidification latent heat per unit time and the amount of hot water.

  First, a melt 21 of a predetermined raw material alloy is produced and stored in the hot water storage container 4 of the melting chamber 1 in FIG. Next, the molten metal 21 is discharged from the hot water nozzle 5 onto the cooling roll 7 in the reduced pressure Ar atmosphere, rapidly cooled by contact with the cooling roll 7, and solidified.

  The time during which the molten alloy 21 is cooled by the cooling roll 7 corresponds to the time from when the alloy comes into contact with the outer peripheral surface of the rotating cooling roll 7 until it leaves, during which the temperature of the alloy suddenly decreases, It becomes a supercooled, that is, almost amorphous, solid. Thereafter, the supercooled alloy leaves the cooling roll 7 and flies in an inert atmosphere. While the alloy is in the form of a ribbon, the temperature is further reduced as a result of the heat being taken away by the atmospheric gas. For example, the pressure of the atmospheric gas is set within the range of 10 kPa to normal pressure. By increasing the atmospheric gas pressure to 30 kPa or higher in order to enhance the secondary cooling effect by the atmospheric gas, a thin ribbon-like quenched alloy having a homogeneous structure can be obtained.

  In addition, the rapid cooling method of the molten alloy used in the present invention is not limited to the above-described single roll method, and a strip casting method which is a rapid cooling method in which the flow rate is not controlled by the nozzle orifice may be used. In the case of the strip casting method, since the nozzle orifice is not used, there is an advantage that the melt supply rate can be easily increased and stabilized by adjusting the amount of pouring. However, atmospheric gas is likely to be caught between the cooling roll and the molten metal, and the cooling rate on the quenching surface side may become nonuniform. In order to solve such a problem, it is necessary to reduce the atmospheric pressure of the space in which the cooling roll is placed to the above-described range and suppress the entrainment of the atmospheric gas. Further, although the production efficiency is lowered, a gas atomizing method may be used.

[Cooling rate]
In order to obtain a quenched alloy having a fine structure in which the average size in the minor axis direction of each phase constituting the crystal structure is 40 nm or more and 2 μm or less, the average cooling rate in the temperature range from 1500 ° C. to 600 ° C. Is preferably in the range of 2 × 10 ⁴ ° C./s or more and 2 × 10 ⁶ ° C./s or less, and particularly preferably 2 × 10 ⁵ ° C./s or more.

For example, when a round nozzle having a nozzle orifice diameter of 0.8 mm is used, the average cooling rate 2 × 10 ⁵ ° C./s is 2 × 10 ⁵ ° C./s at a roll peripheral speed of about 3 m / s and 2 × 10 ⁵ at about 30 m / s. ⁶ ° C./s is obtained.

According to the strip casting method, wide strips can be produced continuously, and the amount of alloy that can be processed per unit time can be dramatically increased. On the other hand, the amount of heat removed from the cooling roll needs to be increased accordingly, and finding an appropriate balance between productivity and the amount of heat removed from the roll is an engineering issue. According to experience, the cooling rate that can be realized by the strip casting method is smaller than the cooling rate by the melt spinning method and is about 10 ³ ° C / s to 4 × 10 ⁵ ° C / s. For example, when producing a strip having a width of 10 mm and a thickness of 90 μm using a cooling roll having an internal water cooling structure having a diameter of 500 mm, a roll peripheral speed at which a cooling speed of 2 × 10 ⁵ ° C./s is obtained is 10 to 15 m / s. Is in range. The number of strips produced at the same time is equal to the number of tundish gates, and the melt feed rate per strip is typically in the range of approximately 1-5 kg / min.

  In order to make the cooling rate as described above, it is preferable to set the hot water conditions and the roll peripheral speed of the cooling roll as follows.

[Preferable range of tapping temperature]
The liquidus temperature of the alloy represented by the above composition formula is approximately 1450 ° C., and the tapping temperature needs to be higher. Since primary crystal iron is generated below the liquidus temperature, if the tapping temperature is lower than that, the molten metal containing coarse primary crystal iron is poured into the cooling roll, and the objective of obtaining a fine structure cannot be achieved. Therefore, the molten metal temperature needs to be sufficiently higher than the liquidus temperature. However, if the temperature becomes excessively high, the amount of heat removal required during the ultra-quenching process increases in proportion to the absolute temperature, which is not preferable. An appropriate molten metal temperature is preferably controlled to a narrow range of about 50 ° C. for these reasons from an engineering viewpoint, and is preferably higher than the liquidus temperature by a temperature of 50 ° C. or more and 150 ° C. or less. The actual measured temperature depends on the measuring equipment, and also depends on the measurement position of the molten metal temperature. In general, however, measuring the molten metal temperature in the tundish or in the melting crucible is the most unsuccessful operation. Considering the possibility that the temperature lowers further from the time of measuring the molten metal temperature until reaching the cooling roll, the lower limit value of the tapping temperature is preferably 50 ° C. or more higher than the liquidus temperature.

[Thickness of quenched alloy]
The thickness of the quenched alloy ribbon is almost inversely proportional to the peripheral speed of the cooling roll, proportional to the amount of tapping per strip per unit time, and inversely proportional to the strip width. Since the most important cooling rate is approximately proportional to the peripheral speed of the cooling roll, the range of the quenched alloy ribbon thickness is almost uniquely determined by the preferable range of the cooling rate. Factors that determine these proportional constants include, for example, the diameter of the cooling roll, the heat removal capability of the cooling roll, etc. The latter depends on the material of the roll, the thickness, the internal water cooling structure, the cooling water flow rate, the cooling water flow speed, and the like. . Therefore, the preferred range of the thickness of the quenched alloy ribbon is in a relatively narrow range determined by the production apparatus, but when the thickness is approximately 20 μm, the upper limit of the average cooling rate is 2 × 10 ⁶ ° C./s and the thickness is 200 μm. In this case, the lower limit value 2 × 10 ⁵ ° C / s of the average cooling rate is obtained.

  By defining the thickness of the quenched alloy ribbon, it is guaranteed that the alloy can be produced at an appropriate cooling rate, and the production of the alloy having the structural characteristics of the present invention is facilitated.

  Here, the effect obtained by using the liquid quenching method will be described in detail with reference to a reference experimental example.

(Reference Experiment Example 1)
[Production process of raw material alloy melt]
A predetermined amount of La, Fe and Si raw materials are blended so that a La (Fe, Si) ₁₃ type compound phase having a composition of La (Fe _0.88 Si _0.12 ) ₁₃ is obtained, and a cast alloy is produced using a high frequency melting furnace. did. The cast alloy (as-cast alloy) obtained at this stage is used as a sample (e). In this specification, a cast alloy that has not been heat-treated may be referred to as an as-cast alloy in order to distinguish it from a quenched alloy that has not been heat-treated (as-spun alloy).

[Rapid cooling process]
Using an experimental machine having the same configuration as that shown in FIG. 1, about 10 g of molten cast alloy was injected from a quartz nozzle having a diameter of 0.8 mm onto a Cu roll rotating at 20 m / second to produce an alloy ribbon. The alloy ribbon (as-spun alloy) obtained here is taken as a sample (a).

[Heat treatment process]
Sample (a) was wrapped in Nb foil, placed in a quartz tube, and subjected to heat treatment at 1000 ° C. for 1 hour while being evacuated with a rotary pump (substantially a vacuum of 10 Pa or less). The obtained quenched alloy is designated as sample (b).

Sample (a) was enclosed in a quartz tube evacuated to 10 ⁻² Pa or less and quenched at 1050 ° C. for 24 hours as sample (c), and quenched at 120 hours for sample (c) d).

About 10 g of sample (e) (casting alloy) was sealed in a quartz tube evacuated to 10 ⁻² Pa or less, and a casting alloy that was heat treated at 1050 ° C. for 1, 24, and 120 hours was used as sample (f), Let it be a sample (g) and a sample (h).

[Evaluation]
The crystal structure of each sample was evaluated by an X-ray diffraction method (XRD method). The powder which grind | pulverized each sample to 150 micrometers or less was used. Cu was used for the target. The scan speed was 4.0 ° / min, the sampling width was 0.02 °, and the measurement range was 20 to 80 °.

  Table 1 summarizes the heat treatment conditions of the obtained samples (a) to (h) and the formation phase in the alloy.

  Moreover, the form and composition distribution of each sample were evaluated using an electron beam microanalyzer (EPMA). A sample for EPMA observation was prepared as follows. An EPMA sample was prepared by impregnating each sample alloy with an epoxy resin and polishing the surface, followed by Au deposition with a thickness of about 20 nm. The acceleration voltage of EPMA was 15 kV. The irradiation current is B.I. E. I. (Reflected electron image) was 1.0 nA.

The magnetic properties (magnetocaloric effect) of each sample were evaluated. A material having a large magnetocaloric effect is preferred for the magnetic refrigeration material. Magnetic entropy change −ΔS _mag is used to evaluate the magnetocaloric effect. In general, the greater the -ΔS _mag , the greater the magnetocaloric effect. Using a high magnetic field VSM (sample vibration magnetometer), a magnetization (M) -temperature (T) curve was measured under an applied magnetic field having a constant intensity from 0T to 1T at intervals of 0.2T. -ΔS _mag was calculated using Equation (1).

FIG. 2 shows the XRD measurement results of samples (a), (b), (c) and (d) prepared using the quenched alloy. Further, FIG. 3 shows the temperature dependence of −ΔS _mag obtained for the sample (c). Further, a reflected electron image (BEI) of the sample (c) by EPMA is shown in FIG.

  For comparison, FIG. 5 shows the XRD measurement results of samples (e), (f), (g) and (h) prepared using a cast alloy. In addition, reflected electron images (BEI) of samples (e) and (h) by EPMA are shown in FIGS. FIG. 6B shows a backscattered electron image of a sample having a heat treatment time of 8 hours.

  The difference in structure between the quenched alloy sample and the cast alloy sample will be described while comparing FIG. 2 and FIG.

As can be seen from FIG. 2, in the quenched alloy, a La (Fe, Si) ₁₃ type compound phase (◯ in the figure) is formed immediately after quenching (sample (a)). In the sample (a), (La, Fe, Si) compound phases (▲ in the figure) and α-Fe phases composed of La, Fe, and Si are also formed. When heat treatment is performed for 1 hour (sample (b)), the (La, Fe, Si) compound phase almost disappears and the α-Fe phase also decreases. Thereafter, when the heat treatment time is lengthened, almost no change is observed except that the intensity of the peak derived from the α-Fe phase is slightly increased. In this case, almost all of the quenched alloy is La ( It can be seen that this is a Fe, Si) ₁₃ type compound phase. Moreover, when the backscattered electron image of the sample (c) shown in FIG. 4 is seen, it has a substantially uniform composition distribution over the whole ribbon except that many Fe exist in the edge part of a ribbon. I understand that.

Moreover, as can be seen from the temperature dependence of the magnetic entropy change −ΔS _mag shown in FIG. 3, the quenched alloy (sample (c)) shows a large magnetic entropy change. -ΔS _mag from 0T to 1T was 7.5 Jkg ⁻¹ K ⁻¹ . Currently, Gd (gadolinium) used in a magnetic refrigeration tester operating near room temperature is about −ΔS _mag = 3 Jkg ⁻¹ K ⁻¹ from 0T to 1T, which is a large magnetic entropy change. It can be seen that In addition, -ΔS _mag obtained for the sample after removing the oxide layer (thickness of about 2 μm) on the surface of the sample (h) prepared using the cast alloy was 19 Jkg ⁻¹ K ⁻¹ . The reason why -ΔS _{mag of} sample (c) is lower than that of sample (h) is considered to be due to the presence or absence of the oxide layer on the surface. Considering the industrial applicability, it is considered that the effect of shortening the heat treatment time, the raw material cost, and the simplification of the pulverization process are larger than the decrease of -ΔS _mag .

On the other hand, when the XRD measurement results of the samples (e) to (h) prepared using the conventional cast alloy shown in FIG. 5 and the reflected electron image shown in FIG. 6 are viewed, the cast alloy (sample (e) ) Has no La (Fe, Si) ₁₃ type compound phase, and as the heat treatment proceeds, the (La, Fe, Si) compound phase and the α-Fe phase gradually decrease, and La (Fe, Si) ) You can see how the _13- type compound phase is formed. Further, as can be seen from a comparison between FIG. 5 and FIG. 2, in the sample (b) obtained by heat-treating the quenched alloy for 1 hour, the (La, Fe, Si) compound phase disappears almost completely, but the casting In the sample (g) obtained by heat-treating the alloy for 24 hours, the (La, Fe, Si) compound phase remains.

Here, in order to explain the formation process of the La (Fe, Si) ₁₃ type compound by the solid phase diffusion reaction, FIG. 25 shows an example of a cast alloy structure. FIG. 25 is a reflected electron image (BE) by EPMA and a composition image by fluorescent X-rays. The distribution of each element can be confirmed by the fluorescent X-ray image shown below, and it can be seen that the LaFeSi liquid phase acts on the α solid solution and a La (Fe, Si) ₁₃ type compound is formed by the peritectic reaction. According to the quantitative analysis using EPMA, the Si concentration in the α-Fe phase was in the range of 1.8 at% to 8.5 at%. In addition, when the amount of La is increased, a La—Si compound is also generated in addition to these phases. According to thermal analysis, the liquidus of this composition is 1675K, the peritectic point is 1575K, and the freezing point is 1511K. The high freezing point is considered to be the cause of difficulty in homogenization. The above is also true for the quenched alloy, and in order to obtain an appropriate final composition, an alloy in which 1.5 to 10 at% Si is dissolved in the α-Fe phase is considered preferable. It is done. If the amount of solid solution of Si exceeds 10 at%, the ratio of Fe and Si in the La (Fe, Si) ₁₃ type compound produced by the peritectic reaction is closer to the Si-rich side, which may impair the magnetic properties. is there.

Thus, it can be seen that by using a quenched alloy, a La (Fe, Si) ₁₃ -based magnetic alloy having a La (Fe, Si) ₁₃ type compound phase as a main phase can be obtained by a short heat treatment.

(Reference Experiment Example 2)
Next, a description will be given of results of experiments for obtaining an optimum heat treatment time for a La—Fe—Si alloy obtained by the liquid quenching method.

[Sample preparation]
Similar to Experimental Example 1 described above, predetermined amounts of La, Fe, and Si raw materials are blended so that a La (Fe, Si) ₁₃ type compound phase having a composition of La (Fe _0.88 Si _0.12 ) ₁₃ is obtained. A casting alloy was produced using a melting furnace. About 10 g of the molten cast alloy obtained was sprayed from a quartz nozzle having a diameter of 0.8 mm onto a Cu roll rotating at 20 m / sec to produce an alloy ribbon. The alloy ribbon (as-spun alloy) obtained here was used as sample (i).

  Samples (j), (k), (l) were obtained by heating the alloy ribbons obtained by heat-treating the sample (i) in an Ar stream at 1050 ° C. for 1, 5, 10, 30, and 60 minutes, respectively. Sample (m) and sample (n) were used.

Further, similarly to the above-described method, a cast alloy having a composition of La (Fe _1-x Si _x ) ₁₃ (x = 0.10, 0.11, 0.12, 0.13, 0.14) is produced. did. An alloy ribbon was produced from each cast alloy according to the method described above. However, here, the rotational speed of the Cu roll was 10 m / sec.

  The obtained alloy ribbons are wrapped in Nb foil and heat-treated at 1050 ° C. for 1 hour in an Ar stream, and the alloy ribbons obtained are sample (o), sample (p), sample (q), and sample (r), respectively. And sample (s).

On the other hand, for comparison, about 10 g of the above-mentioned cast alloy was sealed in a quartz tube evacuated to 10 ⁻² Pa or less and heat-treated at 1050 ° C. for 120 hours, respectively, sample (t) and sample (u) Sample (v), sample (w) and sample (x).

  Table 2 summarizes the compositions and preparation conditions of samples (i) to (s) obtained by the liquid quenching method and samples (t) to (x) obtained by the casting method.

[Evaluation]
Each sample was evaluated in the same manner as in Experimental Example 1. FIG. 7 shows the results of evaluating the crystal structures of sample (i), sample (j), sample (k), sample (l), sample (m) and sample (n) by XRD.

As can be seen from FIG. 7, when the alloy ribbon (sample (i)) immediately after quenching is subjected to heat treatment, the α-Fe phase is reduced even if the heat treatment time is only 1 minute (sample (j)). Observed clearly. From this, it is considered that even in a very short time (for example, 1 second), the effect of increasing the La (Fe, Si) ₁₃ type compound phase can be obtained more than immediately after the rapid cooling.

  Even if the heat treatment time is further increased, the intensity of the peak derived from the α-Fe phase does not change until the heat treatment time is 1 hour (sample (n)). As described above with reference to FIG. 2, when the heat treatment time is 24 hours (sample (c)), the α-Fe phase increases conversely. From the dependence of the diffraction peak intensity derived from the α-Fe phase on the heat treatment time, it is considered that the α-Fe phase does not increase until the heat treatment time is about 1 hour. That is, the optimum heat treatment time for the quenched alloy is about 1 hour or less.

  Next, the results of observing the fracture surfaces of sample (i), sample (k), and sample (n) with a field emission scanning electron microscope (FE-SEM) are shown in FIGS. 8A, 8B, and 8C. Shown in

  As can be seen from FIG. 8 (a), a particulate microstructure of 1 μm or less is observed in the alloy ribbon (sample (i)) immediately after quenching, but in the sample (k) subjected to heat treatment for 5 minutes. As can be seen from FIG. 8B, a relatively large particulate structure of about 1 μm is formed. Furthermore, in the sample (n) in which the heat treatment time is 1 hour, as shown in FIG.

As described above, the heat treatment reduces the α-Fe phase and increases the La (Fe, Si) ₁₃ type compound phase, thereby promoting homogenization of the structure.

  FIG. 9 shows a sample (o), a sample (p), a sample (q), a sample (r), a sample (s), a sample (t), a sample (u), and a sample (v) of the liquid quenching method. The results of evaluating the magnetic properties (magnetocaloric effect) of sample (w) and sample (x) are shown.

And the temperature dependence of the magnetic entropy change -Derutaesu _mag of each sample of the liquid quenching method, comparing the temperature dependence of the magnetic entropy change -Derutaesu _mag of each sample of the corresponding casting, -Derutaesu _mag in any 205K or less The maximum value of 15 is 21-21 Jkg ⁻¹ K ⁻¹ , and when it exceeds 205 K, the maximum value of −ΔS _mag is 9 Jkg ⁻¹ K ⁻¹ or less, which is almost the same value. That is, a La (Fe, Si) ₁₃ -based magnetic alloy material having a temperature dependence of magnetic entropy change −ΔS _mag equivalent to that obtained by heat-treating the quenched alloy for 1 hour and heat-treated for 120 hours. It can be seen that it can be manufactured.

(Reference Experimental Example 3)
[Surface oxidation]
FIG. 10 shows an XRD diagram of the ribbon surface obtained by heat-treating the as-spun ribbon produced at a roll peripheral speed V _s = 20 m / s at 1050 ° C. (1323 K) for 30 minutes, 1 hour, and 2 hours. On the surface of the ribbon, as the heat treatment time becomes longer, the peak intensity of the La (Fe, Si) ₁₃ type compound phase decreases and the peak of the α-Fe phase increases. This is presumably because the rare earth near the surface formed an oxide during the heat treatment and separated.

  From the above results, it became clear that the homogenization heat treatment can be greatly shortened by using the liquid quenching method. On the other hand, it was suggested that the longer the heat treatment time, the lower the magnetic properties due to the oxidation of the ribbon surface.

(Reference Experiment Example 4)
[Relationship between heat treatment time and magnetic entropy change-ΔS _mag ]
In order to generate a La (Fe, Si) ₁₃ type compound phase in the powder metallurgy sintering process, the optimum heat treatment time should be selected in consideration of the composition and the homogeneity of the structure, the surface condition of the ribbon, etc. Is considered necessary. Therefore, the relationship between the homogeneity of the tissue and the change in magnetic entropy −ΔS _mag was examined.

FIG. 11 shows -ΔS _{mag of} a sample obtained by heat-treating an as-spun ribbon manufactured at a roll peripheral speed V _s = 20 m / s at 1050 ° C. (1323 K) for 5 minutes, 10 minutes, 30 minutes, 1 hour, and 24 hours. Indicates. For comparison, FIG. 11 also shows -ΔS _{mag of} a bulk sample (sample produced by a casting method). The bulk sample is a sample for -ΔS _mag evaluation where the surface layer is removed after heat treatment at 1050 ° C. for 120 hours. The sample after heat treatment was used as it was regardless of the surface state of the ribbon.

Comparing the bulk and the entire ribbon, the peak temperature of −ΔS _mag is located on the high temperature side and the peak height is low in the ribbon. As shown in FIG. 9, since La (Fe _1-x Si _x ) ₁₃ exceeds − = 0.12, −ΔS _mag decreases, so that −ΔS _{mag of the} ribbon sample is generally low. There is a high possibility of composition shift in the sample preparation process.

Paying attention to each -ΔS _mag of the ribbon sample, the peak temperature showing the maximum value shifts to the high temperature side from 5 minutes to 30 minutes, the rise of -ΔS _mag is almost the same at 30 minutes or more, and the peak height Is decreasing. The highest -ΔS _mag was -ΔS _mag = 10.2 Jkg ^-1 K ^-1 after 30 minutes of heat treatment. The peak position is most likely due to the difference in Curie temperature _Tc , that is, the difference in the amount of Si in the La (Fe, Si) ₁₃ type compound phase in the ribbon sample. It is considered that the peak height is dominated by the characteristic distribution in the ribbon, that is, the uniformity of element distribution and the volume reduction of the La (Fe, Si) ₁₃ type compound phase due to oxidation of the ribbon surface layer.

From the above results, it was found that -ΔS _mag is also influenced by the homogeneity of the tissue inside the sample. In addition, after 30 minutes, it is considered that the volume reduction of the La (Fe, Si) ₁₃ type compound phase has a great influence.

(Reference Experimental Example 5)
[Relationship roller peripheral velocity V _s and the magnetic entropy change -ΔS _mag]
In the melt spinning method, the roll peripheral speed V _s affects the ribbon thickness. V _s changes the cooling rate of the molten metal, affects the crystal grain size of the as-spun alloy, and affects the sintering time for obtaining the target phase or the subsequent heat treatment time. Therefore, the influence of the roll peripheral speed V _s on −ΔS _mag was investigated. FIG. 12 shows -ΔS _mag of a ribbon sample obtained with a target composition of La (Fe _0.88 Si _0.12 ) ₁₃ and roll peripheral speeds V _s = 5, 10, and 20 m / s. The heat treatment was performed at 1050 ° C. for 30 minutes and 1 hour. Bulk value -ΔS _mag = 18.8 Jkg ⁻¹ K ⁻¹ as a whole is low, and it is considered that the deviation of the Si amount is also affecting. The highest value was a sample prepared under conditions of V _s = 10 m / s, 1050 ° C., and 30 minutes, and −ΔS _mag = 14.2 Jkg ⁻¹ K ⁻¹ was obtained. In any V _s , -ΔS _mag was lowered by heat treatment for 1 hour.

As the roll peripheral speed V _s increases, the ribbon becomes thinner, the surface area of the sample increases, and the characteristic deterioration during the heat treatment becomes remarkable. However, it is considered that as the cooling rate increases, the crystal grain size of the as-spun alloy becomes finer and homogenization is more likely to proceed. It is considered that V _s = 10 m / s was the optimum condition for the ribbon thickness and crystal grain size within the condition range examined this time.

Summarizing the above results, in order to obtain a quenching ribbon suitable for powder metallurgy having a high -ΔSmag, a sample having no composition deviation is prepared, and an optimum roll peripheral speed V _s according to the alloy composition is obtained. It is considered that the heat treatment temperature and the heat treatment time may be selected. In addition, since the observation of the first-order magnetic phase transition is strongly influenced by the compositional non-uniformity and the heterogeneous phase, the optimum condition changes due to the slight compositional deviation and the influence of impurities even with the same blending composition. For example, the target composition was La (Fe _0.88 Si _0.12 ) ₁₃ and a sample was prepared again from dissolution, and an experiment similar to that shown in FIG. 12 was performed. The optimum heat treatment time at this time was 1 hour.

(Reference Experimental Example 6)
[Oxygen concentration]
A rapidly cooled alloy was prepared by changing the roll peripheral speed by the melt spinning method. FIG. 13 shows the relationship between the oxygen concentration (Δ) of the as-spun quenched alloy and the oxygen concentration (□) of the magnetic alloy containing the target phase obtained by heat-treating it and the roll peripheral speed. Further, FIG. 14 shows the amount of change in magnetic entropy of the magnetic alloy.

  From these results, a high magnetic entropy change is obtained when the roll peripheral speed is in the range of 3 m / s to 25 m / s. Therefore, it is advantageous to limit the oxygen concentration range of the as-spun alloy to 0.02 mass% to 0.05 mass%, and the oxygen concentration range of the quenched alloy after heat treatment to a range of about 0.25 mass% to 0.8 mass%. I know that there is.

[Other composition]
An experimental example demonstrates that the above-described effects can be obtained for an alloy having a composition represented by the above-described composition formula as well as the above-described composition. [Table 3] shows the composition, quenching conditions, and heat treatment conditions of the samples ((1) to (19)) used in the experiment.

As can be seen from the XRD data shown in FIG. 15, in any of the quenched alloys (samples (1) to (8)) including RE (Ce, Pr, Nd, Sm, Gd, Dy, Tb, Er) other than La It can also be seen that a compound phase having a NaZn ₁₃ type crystal structure is produced.

  Further, as shown in FIGS. 16 and 17, even when 10 at% of RE other than La is added, the state of magnetization change around 190 K does not add RE other than La (sample (9), see FIG. 18). And no different. In addition, it is thought that the same effect is acquired also about other rare earth elements (Pm, Eu, Gd, Ho, and Tm).

Further, as shown in FIGS. 19 and 20, even in an alloy in which all of “A” in the composition formula is Al, a compound phase having a NaZn ₁₃ type crystal structure is generated, causing a magnetic entropy change.

Furthermore, as shown in FIGS. 21 and 22, even in the alloy containing Co in the above composition formula, a compound phase having a NaZn ₁₃ type crystal structure is generated, and a magnetic entropy change is also observed. In particular, since Co substitutes for Fe in the compound phase, there is an advantage that the Curie temperature can be freely set to around room temperature depending on the amount of addition.

In the XRD data of the sample (16) shown in FIG. 23, the peak attributed to the compound phase having the NaZn ₁₃ type crystal structure is observed, which indicates that crystallization can be performed at 900 ° C. Furthermore, in the XRD data of the samples (17) to (19) in FIG. 23, since a peak attributed to the compound phase having the NaZn ₁₃ type crystal structure is observed, the heat treatment time is “1 second”. It can also be seen that crystallization is possible.

  Next, the steps from pulverizing the obtained quenched alloy to the powder metallurgy method will be described.

[Quenched alloy powder]
The rapidly cooled alloy (raw material alloy) produced as described above is pulverized using a power mill, pin mill, hammer mill, ball mill, attritor, jet mill or the like to form a fine powder.

  The final shape of the magnetic refrigeration material to be produced is a thin plate for constructing a structure like a heat exchanger in which a liquid as a heat medium flows between a number of parallelly arranged thin plates. To increase the heat exchange efficiency A plate-like shape that is as thin as possible is required as long as the mechanical strength can be maintained. For this purpose, the finer the particle size of the powder, the more advantageous. However, if the particle size is too fine, the specific surface area increases and the oxygen concentration increases. Therefore, the lower limit value of the size of the quenched alloy powder in the minor axis direction is preferably 2 μm. On the other hand, the upper limit of the size in the minor axis direction of the quenched alloy powder is set according to the final shape as described above. However, assuming a thickness of about 2 mm, it is difficult to produce a compact before sintering unless the thickness is 200 μm or less. It becomes. From the viewpoint of forming a sintered body having a higher density, the upper limit value of the size in the minor axis direction of the more preferable quenched alloy powder is less than 10 μm.

[Molded body]
The method for obtaining a thin plate-like sintered body is composed of basic steps of a fine pulverization step, a primary forming step, and a sintering step.

  The primary molding process for producing a molded body using the alloy powder obtained as described above is, for example, press molding using a mold with a powder material in which a lubricant or the like is added to and mixed with the alloy powder as necessary. (Compression molding) is performed. In addition, there are a method of preparing a compound in which an alloy powder and a binder are mixed and molding using a molding method such as injection molding, and a molding method such as an MIM method and a grease sheet method. Also in the press molding method, a binder may be used to ensure the mechanical strength of the molded body. When using a binder, it is preferable to provide a binder removal step for removing the binder in the molded body before the sintering step. It is important to select a binder that generates a low molecular weight hydrocarbon gas that easily combines with La as much as possible. In particular, it is possible to remove the binder at a low temperature in order to avoid the formation of carbides due to chemical reaction with La. Those are preferred. When hydrogen gas is used for promoting debinding, it is desirable to finish the debinding process at 700 ° C. or lower and to reduce the pressure in a vacuum furnace for dehydrogenation at 700 ° C. or higher.

[Sintered body]
Sintering step, homogeneous NaZn ₁₃ type compound phase may be set as appropriate conditions so obtained. Since the alloy powder of the present invention has the fine structure as described above, the target phase can be obtained by sintering at a relatively low temperature and for a relatively short time. For example, the target phase can be obtained by sintering (heat treatment) for 10 seconds or more at a temperature of 600 ° C. or higher. Typically, the target phase can be obtained at a temperature of 900 ° C. or more and a sintering time of 1 hour or less. Sintering temperatures exceeding 1320 ° C. and sintering times exceeding 8 hours are not required. The sintering time is preferably 4 hours or less.

  If a known hot pressing method or electric discharge sintering method is used, the target phase can be obtained at a temperature of 600 ° C. or more and a sintering time of 1 hour or less.

According to the present invention, La (Fe, which has a magnetocaloric effect in which the magnetic entropy change (−ΔS _mag ) exceeds 5 JK ⁻¹ kg ⁻¹ when the external magnetic field is changed from 0 T to 1 T using powder metallurgy. Si) _13- based magnetic alloy (magnetic refrigeration material) can be obtained.

Moreover, the regenerative heat exchanger and the magnetic refrigeration apparatus described in Japanese Patent Application Laid-Open No. 2003-028532 can be configured by using the La (Fe, Si) ₁₃ -based magnetic alloy according to the present embodiment.

The La (Fe, Si) ₁₃ -based magnetic alloy according to the present invention is particularly preferably used as a magnetic refrigeration working material. For example, as disclosed in Patent Document 1 and Patent Document 2, it is also suitable as a magnetostrictive material. Can be used.

(Example)
Hereinafter, specific examples of the method for producing a La (Fe, Si) ₁₃ -based magnetic alloy according to the present invention will be described with reference to examples, but the present invention is not limited thereto.

Example 1
A quenched ribbon was produced with the La (Fe _0.88 Si _0.12 ) ₁₃ composition at a peripheral speed of 5 m / s and 15 m / s by a strip casting method. The average thickness was 150 μm (standard deviation 15 μm) and 100 μm (standard deviation 10 μm), respectively. Further, when the crystal structure was observed with powder XRD and SEM, the constituent phases were NaZn ₁₃ type La (Fe, Si) ₁₃ , bcc- (Fe, Si), and La rich region (phase) whose crystal structure was unidentified. The average size in the minor axis direction of each phase was 1.5 μm and 1.1 μm, respectively. These ribbons were coarsely pulverized by a power mill and then made into a powder having an average particle size of 6 μm in nitrogen gas using a jet mill. Furthermore, zinc stearate 0.05 mass% as a lubricant and wax 0.1 mass% as a binder were added and mixed. This was filled in a 10 mm × 40 mm mold, and a molded body having a thickness of 2 mm was produced by press molding (compact molding) using the mold. The obtained compact was debindered at 400 ° C. in an Ar stream in a vacuum sintering furnace, and then sintered at 1250 ° C. (1523 K) in an Ar stream of 50 Pa for 3 hours, at about 95% of the true density, A thin plate sintered body having dimensions of 8.4 mm × 33.5 mm × 1.7 mm was obtained. 50 sintered bodies were produced by this method, and reproducibility in this step was confirmed. Moreover, the magnetic entropy change of the obtained sintered compact was measured. For the samples prepared at the cooling roll peripheral speeds of 5 m / s and 15 m / s, the transition temperature is 190 K and 199 K, and the average ΔSmag between the applied magnetic fields 0T and 1T is −7 Jkg ⁻¹ K ⁻¹ , − It was 9 Jkg ⁻¹ K ⁻¹ . From the above, it was confirmed that the magnetic refrigeration work substance can be produced by the powder metallurgical method using the magnetic alloy material of the present invention as a raw material.

(Example 2)
Next, the results of studying the influence of the difference in the alloy production process and the influence of the composition and the manufacturing method on the sintering process will be described.

  In the La-Fe-Si ternary phase diagram of FIG.

As a production process, a cast alloy (as-cast alloy) is produced by a die casting method, and is pulverized, molded and sintered, and an as-spun alloy (quenched) obtained by a melt spinning method (10 m / s). Alloy strips) were compared with those crushed, molded and sintered. When the crystal structure of the as-spun alloy obtained by the melt spinning method was observed with powder XRD and SEM, the constituent phases were NaZn ₁₃ type La (Fe, Si) ₁₃ and bcc- (Fe, Si), The average size in the minor axis direction of the phase was 50 to 100 nm. In addition, in order to compare the structure of the cast alloy and the quenched ribbon, each heat-treated one was also evaluated. Evaluation was mainly performed by X-ray diffraction. In addition, shaping | molding was made into the pellet (D: 4Mgm < ^-3 >) of about (phi) 3 * 3mm, and it sintered using the thermal analyzer.

  FIG. 26 shows the X-ray diffraction result of the quenched alloy. It can be seen that in the as-spun state, the LaFeSi compound phase increases when the composition is La-rich.

  FIG. 27 shows an evaluation of a constituent phase change when a cast alloy (as-cast alloy) is pulverized and molded and sintered at 1100 ° C. (1373 K) by XRD. Even when an alloy having a target composition is used. , (Α-Fe, Si) phase and LaFeSi compound phase have not disappeared.

On the other hand, FIG. 28 shows a case where the quenched ribbon is treated in the same manner, and it can be seen that a substantially single-phase La (Fe, Si) ₁₃ phase is generated. However, at 1100 ° C. (1373 K), the sintered density was 80% or less of the true density.

FIG. 29 and FIG. 30 show the evaluation results of the sintering treatment with the La-rich composition, respectively. In both cases, a La (Fe, Si) ₁₃ phase and a LaFeSi compound phase are observed after sintering. On the other hand, the sintered density reaches about 85% of the true density even at 1100 ° C. (1373 K), particularly in the rapidly cooled ribbon. -Riching seems to be effective for densification.

FIG. 31B shows the relationship between the density of the sintered body and the sintering temperature when compacts (density: about 4 Mg / m ³ ) obtained from quenched alloys having various compositions are sintered for 120 minutes. The true density of the alloy is 7.2 Mg / m ³ . The alloy composition is indicated by a composition formula (atomic ratio) in the figure. In the figure, the solid mark and the solid line are when the quenched ribbon (as-spun alloy) is used as the raw material, and the white mark and the broken line are when the cast alloy (as-cast alloy) is the raw material. It can be seen from FIG. 31B that the quenched alloy has a denser sintered body than the cast alloy. It can also be seen that a composition having a higher La composition ratio is more advantageous for densification.

(Example 3)
Next, details of target phase formation in the sintering process will be described.

FIG. 32 is a reflected electron image of a cross section of a sintered body obtained by sintering a raw material alloy having a composition of La ₇ Fe ₈₂ Si ₁₁ at 1200 ° C. (1473 K) for 2 hours, and the left side uses a cast alloy. The sintered body, the right side is a sintered body using a quenched alloy. The black part in the figure indicates pores, the dark gray part is the α-Fe phase, the light gray part is the target phase (La (Fe, Si) ₁₃ phase), the white part is the LaFeSi compound phase or La-Si phase or La phase. It is observed that the sintered body using the right quenching alloy has fewer and smaller α-Fe phases than the sintered body using the left casting alloy.

FIG. 33 shows a backscattered electron image (BEI) by EPMA of a cross section of a sintered body obtained by sintering a quenched alloy having a composition of La ₁₁ Fe ₇₆ Si ₁₃ at 1200 ° C. (1473 K) for 2 hours. It is a photograph which shows a composition image. In the sintered body shown in FIG. 33, unlike the structure of the sintered body using the cast alloy in FIG. 32 (left side), the grain boundary of the target phase is a plurality of LaFeSi compound phases, La-Si phases, and La phases. It can be seen that a large number of liquid phases were formed during sintering, and as a result, densification of the sintered body was promoted.

FIG. 34 shows the X-ray diffraction results of a sintered body obtained by sintering a quenched alloy having a La ₉ Fe ₇₈ Si ₁₃ composition at various temperatures for 2 hours each. When sintered at 1270 ° C (1543K), the target phase (La (Fe, Si) ₁₃ phase) is almost single phase, but when sintered at a higher temperature of 1320 ° C (1593K), A LaFeSi compound phase is generated. This indicates that the peritectic temperature of the target phase is partially exceeded, and it can be seen that the sintering temperature is preferably less than 1320 ° C. (1593 K).

Example 4
Next, the relationship between the above results and the magnetocaloric effect is shown.

FIG. 31A shows the relationship between the composition ratio of La and the magnetocaloric effect (−ΔSmag). The raw material alloy is an as-spun alloy obtained by a melt spinning method (10 m / s) with a target composition represented by the composition formula La _x Fe _bal Si ₁₃ (x = 5 to 12, bal represents the remainder). A sintered body was produced in the same manner as in Example 1 except that (quenched alloy ribbon) was used and the sintering temperature was 1220 ° C. (1493 K).

  As can be seen from FIG. 31A, when the La composition ratio is up to about 8 at%, −ΔSmag increases as the La composition ratio increases, and decreases after reaching a peak at about 8 at%. The reason is considered as follows.

When the La composition ratio is up to about 8 at%, the higher the La composition ratio, the smaller the amount of α- (Fe, Si) phase generated in the alloy and the shorter the distance that the element should diffuse during the sintering process. . Therefore, sintering can be performed in a short time, densification (see FIG. 31B), and the proportion of the La (Fe, Si) ₁₃ phase that bears the excellent magnetocaloric effect increases, thus increasing −ΔSmag. On the other hand, when the La composition ratio exceeds 8 at%, the ratio of phases other than the La (Fe, Si) ₁₃ phase (for example, LaFeSi phase) increases, and as a result, -ΔSmag decreases.

  From the above, from the viewpoint of magnetocaloric effect, the composition ratio of La of the magnetic material suitable for the powder metallurgy method is in the range of 6 at% to 11 at%, more preferably 7 at% to 9 at%. Within range.

  Thus, it has been found that when a sintered body is produced by a powder metallurgy process using the quenched alloy according to the present invention as a raw material, a dense sintered body can be obtained at a relatively low temperature in a relatively short time.

According to the present invention, it is possible to produce a La (Fe, Si) ₁₃ -based magnetic alloy material with higher production efficiency than before. Therefore, it is possible to provide a magnetic refrigeration working substance and a magnetostrictive material at a lower price than before, and for example, it is possible to provide a magnetic refrigeration apparatus at a practical cost. Since the magnetic refrigeration apparatus does not use a refrigerant unlike the gas compression type, the magnetic refrigeration apparatus is environmentally friendly and has a feature that high energy conversion efficiency can be obtained by using a permanent magnet material in combination.

(A) is sectional drawing which shows the example of whole structure of the apparatus used for the method of manufacturing the quenching alloy by this invention, (b) is an enlarged view of the part by which rapid solidification is performed. It is a figure which shows the measurement result of XRD of the sample (a), (b), (c) and (d) produced using the quenching alloy. It is a diagram showing temperature dependence of the resulting -Derutaesu _mag for sample (c) was prepared using the rapidly solidified alloy. It is a photograph which shows the reflected-electron image (BEI) by EPMA of the sample (c) produced using the quenching alloy. It is a figure which shows the measurement result of XRD of the samples (e), (f), (g), and (h) produced using the casting alloy. It is a photograph which shows the reflected electron image (BEI) by EPMA of the sample produced using the casting alloy. It is a figure which shows the measurement result of XRD of sample (i), samples (j), (k), (l), (m), and (n) produced using the quenching alloy. It is a photograph which shows the structure of the torn surface by FE-SEM of sample (i), (k), and (n) produced using the quenching alloy. A graph (top) showing the temperature dependence of -ΔS _mag obtained for sample (o), sample (p), sample (q), sample (r) and sample (s) prepared using the alloy ribbon; , (Graph) showing the temperature dependence of -ΔS _mag obtained for sample (t), sample (u), sample (v), sample (w) and sample (x) prepared using a cast alloy is there. It is an XRD figure of the ribbon surface heat-processed for each time. 5 min as-spun ribbons were manufactured by a roll peripheral speed V _s = 20m / _s, 10 minutes, 30 minutes, graph showing the 1 hour and 24 hours heat treated samples -ΔS _mag. Is a graph showing the relationship between -Derutaesu _mag of ribbon samples obtained with the roll peripheral speed. It is a graph which shows the oxygen concentration ((triangle | delta)) of the quenching alloy of as-spun and the oxygen concentration ((square)) of the magnetic alloy containing the target phase obtained by heat-processing this, and a roll peripheral speed. It is a graph which shows the magnetic entropy variation | change_quantity of the magnetic alloy shown in FIG. It is a figure which shows the XRD data of the alloy containing RE other than La. It is a graph which shows the measurement data of the Curie temperature of the alloy containing RE other than La. It is a graph which shows the measurement data of the Curie temperature of the other alloy containing RE other than La. It is a graph which shows the measurement data of the Curie temperature of the alloy which does not contain RE other than La. It is a figure which shows the XRD data of the quenching alloy containing Al. It is a graph which shows-(DELTA) Smag of the quenching alloy containing Al. It is a figure which shows the XRD data of the quenching alloy containing Co. It is a graph which shows-(DELTA) Smag of the quenching alloy containing Co. It is a figure which shows the XRD data of a magnetic alloy with short heat processing time (1 second). It is a ternary phase diagram of La-Fe-Si. It is the photograph which shows the backscattered electron image (BEI) and composition image by EPMA of an alloy sample. It is a figure which shows the measurement result of XRD of a quenching alloy. It is a figure which shows the measurement result of XRD which shows the structural phase change at the time of sintering a casting alloy at 1100 degreeC. It is a figure which shows the measurement result of XRD which shows the structural phase change at the time of sintering a rapidly cooled alloy at 1100 degreeC. It is a figure which shows the measurement result of XRD which shows the structural phase change at the time of sintering a La rich cast alloy at 1100 degreeC. It is a figure which shows the measurement result of XRD which shows the structural phase change at the time of sintering a La rich quenching alloy at 1100 degreeC. It is a graph which shows the relationship between the composition ratio of La, and a magnetocaloric effect (-(DELTA) Smag). It is a graph which shows the relationship between the density of the sintered compact obtained from the quenching alloy of various compositions, and sintering temperature. La ₇ Fe ₈₂ Si ₁₁ are photographs showing the reflection electron image according to the cross section of EPMA of the sintered body obtained by sintering the alloy at 1200 ° C. the composition. La ₁₁ Fe ₇₆ Si ₁₃ is a photograph showing a backscattered electron image (B.E.I) and composition image by EPMA of a cross section of the sintered body obtained by sintering at 1200 ° C. The quenched alloy composition. Is a diagram showing an X ray diffraction results of the obtained sintered body by sintering a rapidly solidified alloy of La ₉ Fe ₇₈ Si ₁₃ compositions at various temperatures.

Explanation of symbols

1b, 2b, 8b, and 9b Atmospheric gas supply port 1a, 2a, 8a, and 9a Gas exhaust port 1 Melting chamber 2 Quenching chamber 3 Melting furnace 4 Hot water storage vessel 5 Hot water discharge nozzle 6 Funnel 7 Rotating cooling roll 21 Molten metal 22 Alloy ribbon

Claims (11)

A method for producing a sintered body of a magnetic alloy, comprising:
Preparing a molten alloy raw material having a predetermined composition;
By rapidly solidifying the molten alloy material under the condition that the average cooling rate in the temperature range from 1500 ° C. to 600 ° C. is in the range of 2 × 10 ⁴ ° C./s to 2 × 10 ⁶ ° C./s , the composition formula _{Fe 100-abc RE a a b} Co c (RE is a rare earth element necessarily including La, a is Si or Al, 6at% ≦ a ≦ 11at %, 8at% ≦ b ≦ 18at%, 0at% ≦ c ≦ 9 at%), and a two-phase structure comprising an α-Fe phase and a (RE, Fe, A) phase containing RE of 30 at% to 90 at%, or α-Fe phase, 30 at% to 90 at% Having a three-phase structure composed of (RE, Fe, A) phase containing RE and RE (Fe, A) ₁₃ compound phase of NaZn ₁₃ type crystal structure, and the average size in the minor axis direction of each phase is 40 nm or more Rapid cooling within the range of 2μm or less Forming gold,
Obtaining a powder comprising particles having a minor axis size in the range of 2 μm to 200 μm by pulverizing the quenched alloy;
Obtaining a molded body by molding the powder;
Sintering the molded body; and
Including
A method for producing a sintered body of a magnetic alloy, wherein the step of sintering the compact is the only heat treatment step for the powder . A in the composition formula is less 7 atomic% or more 9 atomic%, the production method of the sintered body of magnetic alloy of claim 1. Wherein (RE, Fe, A) phase is REFeSi compound phase method for producing a sintered body of magnetic alloy according to claim 1 or 2. Wherein Co is the alpha-Fe phase, the (RE, Fe, A) phase and the RE (Fe, A) ₁₃ is substituted for Fe of at least one phase of the compound phase, claims 1 to 3 The manufacturing method of the sintered compact of the magnetic alloy in any one of. The manufacturing method of the sintered compact of the magnetic alloy in any one of Claim 1 to 4 whose oxygen content rate of the said powder is 0.07 atomic% or more and 0.18 atomic% or less. The method for producing a sintered body of a magnetic alloy according to any one of claims 1 to 5 , wherein the quenched alloy has a thickness of 2 µm to 200 µm. The magnetic alloy according to any one of claims 1 to 6 , wherein in the step of forming the quenched alloy, a tapping temperature of the alloy raw material is higher than a liquidus temperature of the alloy raw material by a temperature of 50 ° C or higher and 150 ° C or lower. The manufacturing method of the sintered compact of this. The method for producing a sintered body of a magnetic alloy according to any one of claims 1 to 7 , wherein in the step of forming the quenched alloy, a roll peripheral speed of the cooling roll is in a range of 3 m / s to 30 m / s. . The method for producing a sintered body of a magnetic alloy according to any one of claims 1 to 8, wherein the size of the particles constituting the powder in the minor axis direction is less than 10 µm. The said sintering process is a manufacturing method of the sintered compact of the magnetic alloy in any one of Claim 1 to 9 including the process of sintering within the temperature range of 600 to 1320 degreeC. The method for producing a sintered body of a magnetic alloy according to claim 10 , wherein in the sintering step, the time for sintering within the temperature range is 10 seconds or more and 8 hours or less.