CN113061796A

CN113061796A - Iron-based ceramic composite material and coating on surface of aluminum alloy and preparation method of iron-based ceramic composite material and coating

Info

Publication number: CN113061796A
Application number: CN202110169137.XA
Authority: CN
Inventors: 于慧君; 迟一鸣; 陈传忠
Original assignee: Shandong University
Current assignee: Shandong University
Priority date: 2021-02-07
Filing date: 2021-02-07
Publication date: 2021-07-02
Anticipated expiration: 2041-02-07
Also published as: CN113061796B

Abstract

The invention belongs to the technical field of preparation of ceramic composite coatings, and particularly relates to an iron-based ceramic composite material on the surface of an aluminum alloy, a coating and a preparation method thereof. The coating material is formed by mixing iron-based self-fluxing alloy powder, ceramic powder and titanium powder, wherein the mass fraction of the ceramic material is 3-20 omega t%, the mass fraction of the Ti powder is 5 omega t% -50 omega t%, and the balance is Fe-based self-fluxing alloy powder. The ceramic material is boron carbide or boron nitride. Surface modification of a substrate by laser alloyingAnd (4) obtaining a uniform and compact coating. The alloyed layer is mainly made of Fe₄Al₁₃、Al₃Ti intermetallic compound and TiB₂And second phases such as TiC and TiN. TiB formed in situ₂Most of them are hexagonal plate structures or short rod structures, which are dispersed or form agglomerates. Most of TiC and TiN are near-spherical or irregular polyhedral small particles distributed in TiB₂Ambient or intermetallic intergranular. The presence of hard ceramic particles and intermetallic compounds in the alloyed layer greatly improves the surface properties, so that the wear resistance of the coating can reach 14 times of that of the substrate.

Description

Iron-based ceramic composite material and coating on surface of aluminum alloy and preparation method of iron-based ceramic composite material and coating

Technical Field

The invention belongs to the technical field of preparation of ceramic composite coatings, and particularly relates to an iron-based ceramic composite material on the surface of an aluminum alloy, a coating and a preparation method thereof.

Background

The information in this background section is only for enhancement of understanding of the general background of the invention and is not necessarily to be construed as an admission or any form of suggestion that this information forms the prior art that is already known to a person of ordinary skill in the art.

Aluminum is the most widely used class of non-ferrous metal materials in modern industries. Aluminum has a low density, about one third of iron; the plasticity is good, the processing is easy, and various section bars and plates can be supported; the chemical property is active, a layer of compact oxide film can be formed in the air, and the coating has certain corrosion resistance; the specific strength is high; the thermal conductivity is good. The aluminum alloy formed by adding certain elements can have higher strength while keeping the advantages of light weight and the like, becomes an ideal structural material, and has wide application prospect in the fields of aerospace, automobile manufacturing, transportation and the like. However, aluminum and its alloys have low surface hardness and poor wear resistance, which severely limits its application in many fields. Therefore, the improvement of the hardness and wear resistance of the surface of the aluminum alloy is a problem to be solved. With the development of high-power lasers, laser surface strengthening technology is rapidly emerging. The laser beam has the characteristics of high energy density, good coherence, good directivity and the like, and the surface of the matrix is rapidly melted and solidified in a very short time after absorbing the laser, so that the nonequilibrium structure transformation is generated, and the surface performance of the matrix is improved. Compared with conventional surface treatment modes such as spraying, vapor deposition, anodic oxidation and the like, the method has the advantages of high bonding strength, small influence on a matrix, wide application range and the like. The surface of the aluminum alloy is subjected to laser surface treatment, so that the hardness and the wear resistance of the aluminum alloy are improved, and the surface of the aluminum alloy becomes a hot spot for surface modification of the aluminum alloy.

Compared with other common laser cladding base materials such as steel, titanium alloy and the like, the laser modification of the surface of the aluminum alloy has more difficulty. Firstly, aluminum alloys have low melting points and are easily burned, and the reflectivity to laser is extremely high. During laser alloying, a proper material system needs to be selected and technological parameters need to be continuously optimized so as to ensure that a coating with excellent quality is obtained. Secondly, the aluminum alloy is easy to be oxidized to form aluminum oxide, and the oxide film has high melting point and high specific gravity, which seriously affects the quality of the alloying layer. In the laser alloying process, the flow of the protective gas must be strictly controlled to prevent the molten pool from being oxidized.

In the aluminum alloy laser surface alloying technology, designing a proper coating is a very critical step. At present, laser alloying coatings mainly comprise three major types of metals and alloys, ceramic materials and metal matrix ceramic composite materials. Among them, the most commonly selected and widely studied coating materials are self-fluxing alloys such as Ni-based and Co-based. The self-fluxing alloy has good wettability with an aluminum matrix, and a dense coating for metallurgical bonding is easy to obtain. The hardness and wear resistance of the surface of the material can be greatly improved after the alloy powder is made into a coating. Fe-based self-fluxing alloys have excellent oxidation resistance and better high temperature performance and are less expensive than Ni-based self-fluxing alloy materials and are therefore considered to be more promising materials. In our previous research, Fe-Al powders with different proportions are designed, and an alloying layer is prepared on the surface of 6061 aluminum alloy. The alloying layer is mainly made of FeAl and Fe₃Al and FeAl₃The intermetallic compounds are distributed on the alpha-Al substrate in a needle shape, a strip shape and a block shape. With the increase of the Fe content in the mixed powder, the hardness and the wear resistance of the alloying layer are obviously improved.

Ceramic materials such as carbides, oxides and borides have good mechanical properties and are often used in laser surface modification. Under the action of laser, ceramic phases with higher melting points may remain in the coating, and the original high strength and high hardness of the ceramic phases are maintained; on the other hand, part of the ceramic particles may decompose under the action of the high energy beam and react with other elements in the coating to form new compounds. The commonly used ceramic materials are TiC, SiC and B₄C、Al₂O₃And the like. In addition, a great deal of research has been conducted to show thatThe direct addition of ceramic particles may increase defects in the coating due to poor wetting, in contrast to ceramic reinforcing phases prepared in an in situ generation manner, which have better compatibility.

The metal-based composite material can have good wettability and toughness of metal and higher hardness and wear resistance of ceramic particles, so that the metal-based ceramic composite coating prepared by laser alloying can obtain better comprehensive performance.

Disclosure of Invention

Aiming at the problems in the prior art, the invention provides a novel iron-based ceramic composite material on the surface of aluminum alloy, a coating and a preparation method. The ceramic-reinforced iron-based composite coating is prepared on the surface of the material, so that the hardness and the wear resistance of the material are greatly improved.

In order to achieve the technical purpose, the invention adopts the following technical scheme:

the iron-based ceramic composite material for the aluminum alloy surface is formed by mixing Fe-based self-fluxing alloy powder, ceramic powder and pure Ti powder, wherein the content of the ceramic material is 3-20 omega t%, the content of the Ti powder is 5-50 omega t%, and the balance is the Fe-based self-fluxing alloy powder, and the ceramic material is boron carbide or boron nitride. The boron nitride is hexagonal boron nitride (h-BN), cubic boron nitride (c-BN), rhombohedral boron nitride (r-BN) or wurtzite boron nitride (w-BN); the chemical composition of the Fe-based self-fluxing alloy powder was Cr 16.5 wt.%, B0.9 wt.%, Si0.8 wt.%, C0.12 wt.%, the balance being Fe and unavoidable impurities. In order to expand the variety of the iron-based alloying layer material, mixing of boron carbide, boron nitride ceramic powder with iron-based powder was attempted. Boron carbide is second only in hardness to diamond and cubic boron nitride and is a common coating strengthening material. Boron nitride also has high hardness and strength, good corrosion resistance and excellent high-temperature performance. The light material with high hardness, high wear resistance and high toughness can be obtained by compounding the boron carbide, the boron nitride and the metal. Ti is a surface active element and has a strong affinity with B, C, N. Ti is added into the coating to reduce the surface tension, and TiB is generated in situ through chemical reaction₂TiC, TiN and other ceramics as second phase strengthening particlesThe coating has good wettability, reduces the tendency of crack generation and further improves the quality of the coating.

In some embodiments, the ceramic material is B₄And when the content of C is 5-18 omega t.%, the content of Ti powder is 10-45 omega t.%, and the balance is Fe-based self-fluxing alloy powder, the coating has high hardness and good wear resistance. Further, B₄The content of C is 8-12 omega t.%, the content of Ti powder is 20-40 omega t.%, and the balance is Fe-based powder, so that the wear resistance is better; further, B₄The content of C is 10 omegat.%, the content of Ti powder is 30 omegat.%, the content of Fe-based powder is 60 omegat.%, and the wear-resisting property of the coating is the best.

In some embodiments, the ceramic material is hexagonal boron nitride, the content of the hexagonal boron nitride is 3-15 wt.%, the content of the Ti powder is 5-25 wt.%, and the balance is Fe-based self-fluxing alloy powder, so that the wear resistance of the coating is good. Furthermore, the content of the hexagonal boron nitride is 3-8 omega t.%, the content of the Ti powder is 12-18 omega t.%, and the balance is Fe-based powder, so that the wear resistance is better; further, the content of hexagonal boron nitride was 5 ω t.%, the content of Ti powder was 15 ω t.%, and the content of Fe-based powder was 80 ω t.%.

A preparation method of an iron-based ceramic composite coating on the surface of an aluminum alloy comprises the following steps:

and polishing the surface of a substrate material to remove oxide skin and clean the substrate material, grinding the Fe-based self-fluxing alloy powder, the Ti powder and the ceramic powder until the Fe-based self-fluxing alloy powder, the Ti powder and the ceramic powder are uniformly mixed, and presetting a binder on the surface of the substrate to perform laser alloying to obtain the laser alloying iron-based ceramic composite coating. The ceramic material is boron carbide or boron nitride, and the substrate is 2 series, 4 series, 5 series, 6 series and 7 series aluminum alloy; further, the base material is 6061 aluminum alloy, and the main phase in the coating comprises Fe₄Al₁₃、Al₃Ti、TiB₂、TiC、TiN、CrB、Cr₂B、Cr₂₃C₆、Al₄C₃And the like. The binder is water glass, and the volume ratio of the water glass to the water is 1: 3.

In some embodiments, the laser alloying power is 1-3 kW, the scanning speed is 9-15 mm/s, the laser spot diameter is 3-4 mm, and the lap joint rate is 30-50%; further, the laser power is 1.7-2.3 kW, the scanning speed is 9-12 mm/s, the spot diameter is 3mm, and when the lap joint rate is 30%, the surface quality of the coating is good; furthermore, the laser power is 2kW, the scanning speed is 12mm/s, the spot diameter is 3mm, the lap joint rate is 30%, and the surface quality of the coating is best.

In some embodiments, in order to prevent the oxidation of the molten pool, argon is used as a protective gas in the laser alloying process, and the gas flow is 10-12L/min; further, the effect is best when the argon flow is 10L/min.

The iron-based ceramic composite coating prepared by the preparation method and the application thereof in the fields of aerospace, mechanical manufacturing, electronics or transportation and the like.

The invention has the beneficial effects that:

(1) according to the invention, boron carbide or boron nitride and pure titanium are mixed with the iron-based alloy, and the surface of the material is modified through laser alloying, so that a uniform and compact alloying layer is obtained. In coating TiB₂The second phase particles of TiC, TiN and the like are fine in size and are dispersed and distributed on the intermetallic compound substrate as a strengthening phase, so that the wear resistance of the coating can be greatly improved, and the service life of the material in a wear environment is prolonged.

(2) The invention adds Ti in the coating to reduce the surface tension and simultaneously reduces the surface tension by adding Ti in the coating and B in the coating₄C. BN generates chemical reaction, the wettability of a ceramic phase and a metal matrix is improved, the tendency of generating cracks is reduced, and the surface quality of the coating is greatly improved.

(3) The method is simple, strong in practicability and easy to popularize.

Drawings

The accompanying drawings, which are incorporated in and constitute a part of this specification, are included to provide a further understanding of the invention, and are incorporated in and constitute a part of this specification, illustrate exemplary embodiments of the invention and together with the description serve to explain the invention and not to limit the invention.

FIG. 1 is a cross-sectional profile of an alloyed layer under different parameters in example 1 of the present invention: (a) p1.7 kW, V15 mm/s, (b) P2 kW, V15 mm/s, (c) P2.3 kW, V12 mm/s, (d) P2 kW, V9 mm/s, (e) P2 kW, V12 mm/s;

FIG. 2 shows the X-ray diffraction results of the alloyed layers of different samples according to example 1 of the invention: (a) sample # 1, (b) sample # 2, (c) sample # 3, (d) sample # 4;

FIG. 3 is a cross-sectional profile of various sample alloying layers in example 1 of the present invention: (a) sample # 1, (b) sample # 2, (c) sample # 3, (d) sample # 4;

FIG. 4 is the microstructure morphology of the 1# sample alloyed layer in example 1 of the present invention: (a) macroscopic, (b) high-macroscopic cluster structure;

FIG. 5 is a back-scattered electron image and a surface composition distribution of the alloyed layer structure of sample No. 1 in example 1 of the present invention;

FIG. 6 is the microstructure morphology of the 2# sample alloyed layer in example 1 of the present invention: (a) low power, (b) high power;

FIG. 7 shows the microstructure of the alloyed layer of sample No. 3 in example 1 of the present invention: (a) low power, (b) high power;

FIG. 8 is a surface composition analysis of the 3# sample alloyed layer in example 1 of the present invention;

FIG. 9 shows TiB in the 3# sample alloying layer in example 1 of the present invention₂The shape of the/TiC composite structure: (a) low power, (b) high power;

FIG. 10 shows TiB in example 1 of the present invention₂Transmission electron microscope bright field image (a-b), diffraction pattern (c-e), high resolution lattice image (g) and Fe of/TiC composite structure₄Al₁₃The diffraction pattern (f) of (a), the high-resolution lattice image (h);

FIG. 11 shows the morphology of the alloyed layer structure of the No. 4 sample in example 1 of the present invention: (a) low power, (b) high power;

FIG. 12 shows a transmission electron microscope bright field image (a), a high resolution lattice image (b) and an electron diffraction pattern (c, d) of the grain structure of the 4# sample alloying layer in example 1 of the present invention;

FIG. 13 is a microhardness profile of an alloyed layer in example 1 of the present invention;

FIG. 14 is a graph of the coefficient of friction of the 6061 substrate and the sample alloyed layer in example 1 of the invention;

fig. 15 is a laser confocal image and cross-sectional profile of the wear scar of the 6061 substrate and the alloyed layer in example 1 of the present invention: (a) matrix, (b) sample # 1, (c) sample # 2, (d) sample # 3, (e) sample # 4;

FIG. 16 shows the wear scar morphology of the 6061 matrix and the 1-4 # sample alloyed layer in example 1 of the present invention: (a) matrix, (b) sample # 1, (c) sample # 2, (d) sample # 3, (e) sample # 4;

FIG. 17 shows the results of X-ray diffraction of alloyed layers of sample No. 5 and sample No. 6 in example 2 of the present invention: (a) sample No. 5; (b) sample No. 6;

FIG. 18 is SEM images of the alloyed layers of sample No. 5 and sample No. 6 at the interface with the substrate in example 2 of the present invention: (a) sample No. 5; (b) sample No. 6;

FIG. 19 is a distribution diagram of the surface composition of a typical structure in the middle of an alloyed layer of No. 5 sample in example 2 of the present invention;

FIG. 20 is an SEM image of the precipitated phase in the 5# sample alloyed layer in example 2 of the present invention and the corresponding EDS analysis: (a) morphology phase; (b) point 1; (c) point 2; (d) point 3;

FIG. 21 is a graph showing a distribution of near-spherical structures and Al, Ti, and N elements in an alloyed layer in example 2 of the present invention: (a) morphology phase; (b) element line distribution;

FIG. 22 is an SEM image of peritectic structures in the alloyed layer of sample No. 6 in example 2 of the present invention: (a) a bottom; (b) a middle part;

FIG. 23 is a plot of microhardness versus distance for the alloyed layers of sample No. 5 and sample No. 6 in example 2 of the present invention;

FIG. 24 shows the wear profiles of the alloyed layers of the aluminum alloy substrate, the 5# sample and the 6# sample in example 2 of the present invention: (a) a substrate; (b) sample No. 5; (c) sample No. 6;

fig. 25 is a laser confocal image of a wear scar and a cross-sectional profile of an alloyed layer in example 2 of the present invention: (a) sample # 5, (b) sample # 6.

Detailed Description

It is to be understood that the following detailed description is exemplary and is intended to provide further explanation of the invention as claimed. Unless defined otherwise, all technical and scientific terms used herein have the same meaning as commonly understood by one of ordinary skill in the art to which this invention belongs.

It is noted that the terminology used herein is for the purpose of describing particular embodiments only and is not intended to be limiting of exemplary embodiments according to the invention. As used herein, the singular forms "a", "an" and "the" are intended to include the plural forms as well, and it should be understood that when the terms "comprises" and/or "comprising" are used in this specification, they specify the presence of stated features, steps, operations, devices, components, and/or combinations thereof, unless the context clearly indicates otherwise.

The present invention is described in further detail below with reference to specific examples, which are intended to be illustrative of the invention and not limiting.

Example 1:

1 Material design and Process

The material used in this example was 6061 aluminum alloy. The chemical composition and main properties are shown in tables 1 and 2. The coating material is JG3 iron-based self-fluxing alloy powder (more than or equal to 99.5 omega t.%, 325 meshes), B₄C (not less than 99.5 omega t.%, 50-100 mu m) and pure Ti powder (not less than 99.5 omega t.%, 325 meshes) are mixed according to a certain proportion. The composition of the JG3 iron-based self-fluxing alloy is shown in table 3. 6061 aluminum alloy plates were cut into test pieces of 20mm × 10mm × 10mm and 20mm × 20mm × 10mm using a DK77-40 model electric discharge numerically controlled wire cutting machine, which was used for wear tests. And (3) polishing the oxide film on the surface of the sample by using sand paper, cleaning the oxide film in an alcohol solution, and airing the oxide film for later use. And weighing alloying powder with a certain mass ratio by using an electronic balance, and manually grinding the alloying powder in a mortar uniformly. Using water glass (Na)₂O·nSiO₂:H₂O1: 3, vol.%) was used as a binder, and the uniformly mixed powder was pre-coated on the surface of the sample to control the thickness to be about 0.6 mm.

TABLE 16061 chemical composition of aluminum alloy matrix (mass fraction,. omega.t.%)

Table 2 chemical composition of Fe-based alloy powder (mass fraction,. omega.t.%)

The sample was scanned using a YLS-4000 fiber laser. In the preliminary test, the addition amounts of B4C and Ti powder are both 10 ω t.%, different laser process parameters are designed, the surface quality of the obtained coating is shown in Table 3, and the section morphology is shown in FIG. 1. When the laser power is 2kW and the scanning speed is 12mm/s, the obtained coating has a flat and smooth surface, no obvious defects in the coating and the best coating quality. Therefore, the laser process parameters of P2 kW, V12 mm/s and the laser spot diameter of 3mm are selected in the subsequent experiments. In order to prevent the oxidation of the molten pool, argon is used as a protective gas in the scanning process, and the gas flow is 10L/min.

TABLE 3 surface quality of the coatings under different process parameters

To investigate the effect of different ceramic contents on the coating, in previous experiments B4C powder was added in mass fractions of 5 ω t.%, 10 ω t.%, 15 ω t.% and 20 ω t.%, respectively. The comparison shows that when the addition amount of the B4C is 5 omega t to 10 omega t, the surface of the coating is smooth, the coating is well combined with a substrate, no obvious cracks and pores exist in the coating, and the wear resistance of the coating is better when the addition amount is 10 omega t. When the addition amount is 15 ω t.% and 20 ω t.%, the hardness of the coating layer is further increased, but the excessive ceramic phase deteriorates the wettability of the coating layer, and more cracks and pores occur. The increase in brittleness is also not favorable for the improvement of wear resistance. Therefore, we chose to add B4C at 10 wt.% in subsequent experiments. As shown in table 4, the coatings with four component ratios were designed according to the difference in Ti powder content in the pre-set coating for laser alloying to obtain an alloyed layer. The alloyed layer was subjected to phase analysis using an UltimaIV X-ray diffractometer from Rigaku. The scanning voltage is 40kV, the current is 40mA, and the scanning speed is 10 DEG/min. The structure morphology of the alloying layer is observed by adopting an S-3400N type scanning electron microscope and a JSM-7800F type field emission high temperature Scanning Electron Microscope (SEM), and the scanning electron microscope is provided with an energy spectrum analysis accessory (EDS) and can be used for component analysis of the structure. The phase was further identified by JEM-2100 transmission electron microscopy.

Table 4 laser alloying sample number and material ratio (mass fraction,. omega.t.%)

The microhardness of the alloyed layer is tested by an HDV-1000 type microhardness tester, the load is 200g, and the loading time is 10 s. Microhardness values were measured every 0.1mm from the surface of the coating to the substrate along the direction of maximum penetration of the coating to analyze the hardness profile, and three measurements were taken at each location and averaged. An HT-1000 type abrasion tester is used for carrying out an abrasion test on an alloying layer, a GCr15 steel ball is adopted as a grinding ball, the diameter of the grinding ball is 4mm, the rotating speed is 280r/min, the load is 750g, the friction radius is 3mm, and the abrasion time is 20 min. And observing the three-dimensional appearance of the grinding mark of the sample by using an LSM800 laser confocal microscope, and obtaining a section profile. The corresponding wear volume is calculated from the cross-sectional profile.

2 results and discussion

2.1 phase composition

FIG. 2 is an X-ray diffraction pattern of an alloyed layer of different samples, and the content of Ti has a significant influence on the phase composition of the alloyed layer.

Samples

1 and 2 had a low Ti content and the alloyed layer consisted primarily of Fe₄Al₁₃、TiB₂、CrB、Cr₂B、Cr₂₃C₆、Al₄C₃And alpha-Al. Under the action of high-energy laser beam, the pre-coating and a thin layer on the surface of the substrate are simultaneously melted to form a molten pool, and a series of complex chemical reactions can occur in the molten pool to generate a plurality of phases. First, B in the molten pool₄C is decomposed into B atoms and C atoms, and since Ti atoms have a greater affinity for B atoms than for C atoms, TiB₂Preferentially, excess B also reacts with Cr in the bath to form Cr-B compounds. With TiB₂Continuously absorbing B atoms to grow, and gradually forming a C-rich B-poor area at the front edge of a solid-liquid interface to further promote the nucleation growth of carbide. Further, Fe is generated₄Al₁₃The Gibbs free energy is lower and can also be formed in a spontaneous reaction in a molten pool. When the addition of Ti was 30 wt.%, the presence of TiC was detected in the alloyed layer. When the Ti content is increased to 45 ω t.%, in addition to TiC, excess Ti atoms in the bath react with Al atoms to form Al₃Ti。

2.2 microstructure

FIG. 3 is the structural morphology of the cross section of the alloyed layer of different samples, and it can be seen that the alloyed layer and the substrate are tightly bonded, no obvious boundary line exists, and the characteristics of good metallurgical bonding are presented. As shown in fig. 3(a), the alloying layer can be roughly divided into three parts from the outside to the inside: an alloying zone, an interfacial zone and a matrix. The thickness of the alloying layer is about 0.8-1 mm, the structure is uniform and compact, and no crack or air hole is found.

Fig. 4 is a typical structure morphology of the alloyed layer of sample # 1, and a large amount of granular, needle-like, short rod-like, cluster-like precipitated phases are observed to be uniformly distributed on the strip-shaped substrate. EDS analysis was performed on the cluster structures (points a and E), the precipitated phases (points B and C, F), and the long substrate (Point D), respectively, and the results are shown in table 4. The results show that: the cluster-like structure (Point A, E) contains mainly B, Ti two elements, possibly TiB₂(ii) a The irregular granular (Point C) and acicular (Point B) textures contain B, C and Cr, possibly Cr-B or Cr-C; the strip-shaped substrate (Point D) mainly contains two elements of Fe and Al, and the atomic ratio thereof is about 1: 3.2, possibly Fe₄Al₁₃An intermetallic compound; the short rod-shaped structure (Point F) contains mainly Al and C, possibly Al₄C₃. To further investigate the microstructure of the agglomerates, we observed under high power electron microscopy as shown in fig. 4 (b). It can be seen that the interior of the agglomerates is composed of a large number of approximately spherical and polygonal particles, some of which are represented by regular hexagons, such as hexagonal plate-like crystals with a size of about 0.5 μm and a thickness of about 0.2 μm in the figure. TiB₂Has a hexagonal close-packed structure, and TiB is in equilibrium condition according to the crystal growth theory₂The particle morphology is represented by {0001} and

hexagonal prism with wrapped crystal face. Due to the fact that

Small surface atom density and high growth rate, so TiB₂Tending to become flat hexagonal plate-like crystals as shown in the figure. The research shows that TiB₂The reinforcing phase takes on different shapes in different matrices. In the aluminum matrix, TiB₂Typically exhibiting hexagonal platelet structures with dimensions on the order of tens of nanometers to a few micrometers. In addition, quadrangular or hexagonal TiB₂Often tightly connected together by faces or corners to form agglomerates. FIG. 5 is a plot of the backscattered electron image and corresponding areal composition of the alloyed layer structure of sample # 1, wherein the lumpy texture is mainly composed of Ti and B, the irregular granular and acicular texture is rich in Cr elements, the short rod texture is rich in Al elements, and the substrate is rich in Al and Fe elements, further confirming the above analysis.

TABLE 4 EDS analysis of typical structures in alloyed layers

FIG. 6 is the texture of the 2# sample alloyed layer. Similar to sample # 1, the alloyed layer consisted primarily of TiB₂Agglomerates, granular or needle-shaped carbide, boride and a long-strip Fe-Al intermetallic compound substrate.

Fig. 7 and 8 are the structural morphology and the surface composition analysis, respectively, of the upper part of the 3# sample alloyed layer. The upper part of the alloying layer mainly comprises a Fe-Al intermetallic compound substrate and a cluster-shaped structure TiB₂And a granular precipitated phase TiC. TiB in clustered tissue₂The particles are mainly hexagonal prism-shaped or hexagonal sheet-shaped, the size is about 0.5 mu m to 1 mu m, and the TiC generated in situ is mainly irregular polyhedron-shaped or nearly sphere-shapedAnd the size is about 1 μm. Due to TiB₂While precipitating, Ti, C and other atoms are enriched around the TiC particles₂The cluster-like structure is gathered around. In addition, TiB was observed below the alloyed layer₂the/TiC complex structure is shown in FIG. 9. During the solidification process, TiB₂Hexagonal flaky shape is preferentially precipitated and grown, Ti and C atoms are enriched on the surface of the hexagonal flaky shape, and TiC can be precipitated on TiB after a certain nucleation condition is met₂The surface heterogeneous nucleates and grows to form the complex structure shown in fig. 9.

The composite tissue was further observed and analyzed by transmission electron microscopy. FIG. 10(a-b) is a transmission electron microscope bright field image of the complex structure in the alloyed layer, and FIG. (c) is an electron diffraction pattern at the complex structure interface (Zone 1). It can be seen that two sets of electron diffraction spots, labeled SAED 1 and SAED 2, are present at the interface at the same time. The two sets of diffraction patterns were calibrated and the results are shown in FIGS. (d), (e), which confirmed that SAED 1 and SAED 2 were TiC

And TiB₂

Diffraction spots of the ribbon axis. Panel (g) is TiB₂High-resolution lattice image at/TiC composite tissue interface and TiC (111) and TiB on two sides of interface₂(0001) The atomic arrangement of the crystal plane. It can be seen that there is a clear boundary between the two phases and no interfacial reaction product, indicating good bonding between the two. Fig. 10(f) shows the electron diffraction pattern of zone 2 in (b), and fig. (h) shows a high-resolution lattice image of this zone. They can be confirmed to belong to monoclinic system Fe respectively₄Al₁₃Of crystals

Ribbon axis and (022) crystal plane.

FIG. 11 is the texture of the 4# sample alloyed layer. A large amount of in-situ generated ceramic phase is dispersed and distributed among the crystal grains of the dendritic intermetallic compound in a fine granular shape. Addition of more Ti to meltThe relative content of Fe in the pool is reduced, so that Fe in the alloyed layer is reduced₄Al₁₃The content was reduced and the structure became coarse. Al is also formed in the alloyed layer₃The Ti compound has a crystal structure of a DO22 type face-centered tetragonal structure, is low in symmetry and less in slip system, is a phase with high brittleness and poor toughness and tends to have adverse effects on the performance. In the alloyed layer, TiC remains nearly spherical or irregular polyhedral, while TiB₂It is mostly regular hexagonal platelets. FIG. 12(a) is a TEM image of the irregular grain-reinforced phase in the alloyed layer and its surrounding matrix structure, normalized to belong to TiC crystals

Of ribbon axes and alpha-Al

A ribbon axis. FIG. 12(b) is a high resolution lattice image at the interface between two phases, TiC on both sides of the interface

Crystal plane and alpha-Al (200) crystal plane.

2.3 microhardness and abrasion resistance

FIG. 13 is a graph showing the distribution of microhardness of the alloyed layer of samples # 1 to # 4 as a function of distance. The average hardness of the alloyed layers was 356HV_0.2、378HV_0.2、520HV_0.2And 442HV_0.2Is a substrate (70 HV)_0.2)5 to 7 times of the total weight of the composition. The hardness distribution of the alloyed layer of the sample No. 1 and the sample No. 2 is relatively uniform, and the dispersed carbide and boride are main strengthening phases of the coating. When more Ti is added, TiB₂And the amount of TiC increased, the hardness of the coatings of the 3# and 4# samples further increased. Hardness of 3# sample reaches 520HV_0.2. The hardness of the 4# sample is slightly lower than that of the 3# sample, and TiB is dispersed and distributed among grains₂And TiC is the main strengthening phase, with a hardness of about 442HV_0.2. The hardness of the surface of the alloyed layer is slightly lower than that of the middle and bottom, which may be caused by the dilution effect of the matrix Al on the alloyed layer. Microhardness of the substrate in the region near the interfaceThe degree will be slightly higher than that of the 6061 matrix, and this transition region is often referred to as the heat affected zone, with the differential microhardness caused by the rapid heating and cooling action of the laser melt pool.

The coating and substrate were subjected to a dry friction wear test at room temperature and the results of the coefficient of friction test are shown in FIG. 14. At the beginning of wear, the coefficients of friction of both the substrate and the coating rise rapidly, and after a certain period of "running in", both enter a stable friction phase. At this time, the coefficient of friction of the substrate was about 0.46, and the coefficients of friction of the 1-4 # sample coatings were 0.38, 0.32, 0.25, and 0.31, respectively. Therefore, the friction coefficient of the coating is far smaller than that of the matrix, and the wear resistance is better. In previous researches, the inventor has found that in the process of carrying out counter-grinding with a GCr15 steel ball, metal oxides generated by oxidation reaction of Fe, Al and the like in an alloying layer show a certain lubricating effect during dry friction and wear, so that the coating has a certain antifriction effect while resisting wear.

After 20 minutes of abrasion test, the abrasion volumes of the matrix and the coatings of the 1-4 # samples can be calculated to be 2482.5 multiplied by 10 respectively according to laser confocal data in figure 15^-12m³,384.5×10^-12m³,279.5×10^-12m³,179.6×10^- ¹²m³And 266.5 × 10^-12m³The abrasion resistance of the alloying layer is greatly improved. The 3# sample alloyed layer has the best wear resistance, about 14 times that of the substrate. Compared with a matrix, the wear resistance of other sample alloying layers is also improved by 6-9 times. To further analyze the wear mechanism, we observed the wear scar morphology of the alloyed layer and the substrate, as shown in fig. 16. Because the surface hardness of the 6061 aluminum matrix is low, the GCr15 steel ball is easily pressed into the matrix surface during the abrasion process, and a deeper groove and a wider grinding trace are formed (a in figure 16). In addition, the material of the substrate surface is subject to plowing and repeated plastic deformation, gradually breaking, flaking or migrating to the ball surface, causing substantial wear of the substrate surface material, exhibiting typical adhesive wear characteristics, which result is more consistent with loss of wear of the substrate material. In contrast, the alloyed layer is passing throughThe abraded surface was more even and smooth, with only shallow parallel furrows and slight flaking observed, and was characterized to some extent by abrasive wear. The surface hardness of the 3# and 4# samples is relatively higher, and the positions of the grinding balls pressed into the surface are shallower in the abrasion process, so that the width of the grinding marks is smaller, and the abrasion resistance is better. In general, the existence of the hard ceramic particles such as carbide and boride generated in situ and the intermetallic compound substrate greatly improves the wear resistance of the coating.

Example 2

1 Material design and Process

The coating material is JG3 iron-based self-fluxing alloy powder (not less than 99.5 omega t.%, 325 meshes), h-BN (not less than 99.5 omega t.%, 50-100 mu m) and pure Ti powder (not less than 99.5 omega t.%, 325 meshes), and is mixed according to a certain proportion. 6061 aluminum alloy plates were cut into test pieces of 20mm × 10mm × 10mm and 20mm × 20mm × 10mm using a DK77-40 model electric discharge numerically controlled wire cutting machine, which was used for wear tests. And (3) polishing the oxide film on the surface of the sample by using sand paper, cleaning the oxide film in an alcohol solution, and airing the oxide film for later use. And weighing alloying powder with a certain mass ratio by using an electronic balance, and manually grinding the alloying powder in a mortar uniformly. Using water glass (Na)₂O·nSiO₂:H₂O1: 3, vol.%) was used as a binder, and the uniformly mixed powder was pre-coated on the surface of the sample to control the thickness to be about 0.6 mm. The sample was scanned using a YLS-4000 fiber laser. In the early test, the addition amounts of h-BN and Ti are respectively 5 omegat.% and 10 omegat.%, the laser power P is 1.7-2.3 kW, the scanning speed V is 9-15 mm/s, the spot diameter is 2-4 mm, and the lap joint rate is 30-50%. It was found that too high laser power or too low scanning speed resulted in too high laser energy density of the coating and severe burning of the coating and the aluminum alloy substrate. On the contrary, when the laser power is too low or the scanning speed is too high, the laser energy density is too low, so that the coating material is not sufficiently melted, the bonding strength is reduced, and the coating is easy to fall off. Therefore, P2 kW, V12 mm/s, spot diameter 3mm, and 30% overlap were finally selected for further testing. In order to prevent the oxidation of the molten pool, argon is used as a protective gas in the scanning process, and the gas flow is 10L/min.

In order to investigate the influence of different contents of h-BN and Ti on the coating, coating materials with different proportions are designed and tested, and the sample numbers are shown in Table 5. The 1# to 3# samples investigated the effect of h-BN content. Due to the high melting point and poor wettability of the h-BN ceramic particles, if the addition amount is too high in the laser alloying process, obvious pores and cracks appear on the coating. Therefore, we finally chose to add 5 ω t.% of h-BN. The samples No. 4 to No. 6 are added with Ti with different mass fractions on the basis of h-BN. Ti is used as an active element, can reduce the surface tension and effectively improve the wettability of the coating. In addition, Ti can react with h-BN in the coating material to generate some hard ceramic phases in situ, and the hardness and the wear resistance of the coating can be further improved.

Table 5 alloying layer material ratio (mass fraction,% ω t, balance Fe-based alloy) and surface quality

The alloyed layer was subjected to phase analysis using an UltimaIV X-ray diffractometer from Rigaku. The scanning voltage is 40kV, the current is 40mA, and the scanning speed is 10 DEG/min. The structure morphology of the alloying layer is observed by adopting an S-3400N type scanning electron microscope and a JSM-7800F type field emission high temperature Scanning Electron Microscope (SEM), and the scanning electron microscope is provided with an energy spectrum analysis accessory (EDS) and can be used for component analysis of the structure. The microhardness of the alloyed layer is tested by an HDV-1000 type microhardness tester, the load is 200g, and the loading time is 10 s. Microhardness values were measured every 0.1mm from the surface of the coating to the substrate along the direction of maximum penetration of the coating to analyze the hardness profile, and three measurements were taken at each location and averaged. An HT-1000 type abrasion tester is used for carrying out an abrasion test on an alloying layer, a GCr15 steel ball is adopted as a grinding ball, the diameter of the grinding ball is 4mm, the rotating speed is 280r/min, the load is 750g, the friction radius is 3mm, and the abrasion time is 20 min. And observing the three-dimensional appearance of the grinding mark of the sample by using an LSM800 laser confocal microscope, and obtaining a section profile. The corresponding wear volume is calculated from the cross-sectional profile.

2 phase composition and texture analysis of alloyed layers

Fig. 17 shows XRD diffraction results of alloyed layers of samples # 5 and # 6, respectively. It can be seen that the addition of 15 ω t.% and 20 ω t.% Ti has no significant effect on the composition of the phase of the alloyed layer. The alloyed layers are mainly made of Fe₄Al₁₃、TiN、AlN、TiB₂、Cr₂B、Cr₇C₃、Al₃Ti and alpha-Al phases. It can be seen that, when Ti is added, TiN, AlN and TiB are generated in the alloying layer in situ₂In addition, Ti element can react with Al in the matrix to generate intermetallic compound Al₃Ti。

Fig. 18 is SEM images of the 5# sample and 6# sample, respectively, at the alloyed layer-to-substrate interface. It can be seen that the alloyed layer exhibited a good metallurgical bond with the substrate, and no cracks or pores were found. Furthermore, comparing (a) and (b) in fig. 18, it was found that the dendritic structure at the bottom of the alloyed layer of the sample # 5 was more dense.

FIG. 19 is a typical topographical image of the middle of the alloyed layer of sample # 5 and the corresponding elemental surface distribution map. A plurality of rod-shaped or irregular granular second phases are dispersed in the alloying layer, and the second phases contain a large amount of Ti, N and B elements, which can be TiN or TiB according to the conjecture of XRD₂. The substrate of the alloyed layer mainly contains Fe element, Al element and Cr element, and the substrate of the alloyed layer is Fe as presumed by XRD₄Al₁₃The intermetallic compound and Cr as an alloying element are dissolved in the intermetallic compound in a solid solution, and act as solid solution strengthening to some extent.

Fig. 20(a) is a high magnification view of the particulate second phase in fig. 19. The appearance of the ceramic particles generated in situ can be more clearly seen to be flaky and irregular spherical. EDS analysis was performed on the short rod-like structure, the polygonal plate-like structure and the nearly spherical structure in (a) of FIG. 20, respectively, and the results are shown in (b-d) of FIG. 20. Based on XRD analysis, the flaky tissue or the rod-shaped tissue is presumed to be TiB₂The near-spherical structure is rich in Al, N and Ti elements, and may be a composite structure of AlN and TiN. To further study this complex tissueThe results of line component analysis are shown in FIG. 21. It can be seen that the white area in the center of the grain is high in the content of Ti element and N element, while the gray area around it is high in the content of Al element and N element. Therefore, we consider this complex structure to be a peritectic structure of TiN and AlN in actuality. In the solidification process of the molten pool, because elements such as Ti, N, B and the like have stronger affinity, TiB₂TiN is preferentially precipitated in hexagonal flakes or small particles in the molten pool. Review of the literature reveals that TiB₂During the growth process, due to anisotropy, the material can grow preferentially along a certain direction and finally grow into a short rod shape or a hexagonal prism shape. The partially melted Al atoms in the matrix enter the molten pool, and when a certain condition is satisfied, the molten Al and the TiN particles precipitated first undergo a peritectic reaction to generate a peritectic structure as shown in fig. 21. Similarly, a similar peritectic structure was found in the middle of the alloyed layer of sample No. 6, as shown in fig. 22. In the alloying layer, the ceramic phase generated in situ has good wettability with the matrix, and the defects caused by the difference of physical and chemical properties are reduced. And the in-situ generated second phase structure is fine, the particle diameter is about 1 mu m, and a better strengthening effect can be achieved.

3 hardness and wear resistance of alloyed layer

FIG. 23 is a microhardness profile of the alloyed layers of the 5# sample and the 6# sample. The results show that the average hardness of the alloyed layers is increased to 470HV_0.2And 420HV_0.2And the hardness of the aluminum alloy matrix is about 6 to 7 times of the hardness of the aluminum alloy matrix. The ceramic particles generated in situ and the dense intermetallic compound substrate in the alloying layer are proved to be helpful for improving the performance of the coating. When the content of Ti added was 15 ω t.%, the hardness of the alloyed layer of the 5# sample was higher than that of the 6# sample, indicating that the more Ti was not added, the better. Excessive Ti can cause the structure of intermetallic compounds in the alloying layer to be coarsened, thereby influencing the improvement of the coating performance.

After 20 minutes of abrasion test, the abrasion volumes of the 5# sample and the 6# sample alloyed layer were calculated to be 207.1 × 10 respectively from the laser confocal data in fig. 25^-12m³,340.0×10^-12m³Approximately the wear volume of the aluminum alloy matrix under the same conditions (2482.5X 10)^-12m³) 8.3% and 13.7%. The average friction coefficients of the coating and the substrate are 0.38, 0.41 and 0.46 respectively, which shows that the wear resistance of the alloyed layer is obviously improved compared with the wear resistance of the aluminum alloy substrate from the aspects of abrasion and friction. The profile of the wear scar is shown in FIG. 24. The abrasion mark of the substrate is deep, the surface of the substrate has obvious furrowing and a large amount of plastic deformation, and because the surface hardness of the aluminum alloy is low, a friction pair is very easy to press into the surface of the substrate in the abrasion process and generates plastic deformation under the action of alternating stress, and a large amount of materials are peeled off in severe cases to show the characteristic of adhesive abrasion. In contrast, the wear surface of the alloyed layer was relatively flat with only shallow furrows and a few minor flaking, characterizing significant abrasive wear. The wear of the alloyed layer of sample # 5 was the slightest, with only a small amount of furrowing observed, consistent with the highest hardness and least wear volume results of sample # 5 above. It is worth noting that the improvement of the hardness is beneficial to the improvement of the wear resistance to a certain extent, but the wear resistance is also related to the toughness, the residual stress and the like of the material, and is a more complex comprehensive performance.

It should be noted that the above-mentioned embodiments are only preferred embodiments of the present invention, and the present invention is not limited thereto, and although the present invention has been described in detail with reference to the foregoing embodiments, it will be apparent to those skilled in the art that modifications and equivalents can be made in the technical solutions described in the foregoing embodiments, or equivalents thereof. Any modification, equivalent replacement, or improvement made within the spirit and principle of the present invention should be included in the protection scope of the present invention. Although the present invention has been described with reference to the specific embodiments, it should be understood by those skilled in the art that various changes and modifications may be made without departing from the spirit and scope of the invention.

Claims

1. An iron-based ceramic composite material on the surface of aluminum alloy is characterized in that: the self-fluxing alloy material is formed by mixing Fe-based self-fluxing alloy powder, ceramic powder and pure Ti powder, wherein the content of the ceramic material is 3-20 omega t.%, the content of the Ti powder is 5-50 omega t.%, and the balance is the Fe-based self-fluxing alloy powder, and the ceramic material is boron carbide or boron nitride.

2. The iron-based ceramic composite material with an aluminum alloy surface according to claim 1, wherein: the ceramic material is B₄C, the content of the C is 5-18 omega t.%, the content of Ti powder is 10-45 omega t.%, and the balance is Fe-based self-fluxing alloy powder;

further, B₄The content of C is 8-12 omega t.%, the content of Ti powder is 20-40 omega t.%, and the balance is Fe-based self-fluxing alloy powder;

further, B₄The content of C is 10 omegas, the content of Ti powder is 30 omegas, and the balance is Fe-based self-fluxing alloy powder.

3. The iron-based ceramic composite material with an aluminum alloy surface according to claim 1, wherein: the ceramic material is boron nitride, the content of the boron nitride is 3-15 omegat.%, the content of Ti powder is 5-25 omegat.%, and the balance is Fe-based self-fluxing alloy powder;

further, the content of boron nitride is 3-8 omegat.%, the content of Ti powder is 12-18 omegat.%, and the balance is Fe-based self-fluxing alloy powder;

further, the boron nitride content was 5 ω t.%, the Ti powder 15 ω t.%, and the balance Fe-based self-fluxing alloy powder.

4. The iron-based ceramic composite material with an aluminum alloy surface according to claim 1, wherein: the boron nitride is hexagonal boron nitride h-BN, cubic boron nitride c-BN, rhombohedral boron nitride r-BN or wurtzite boron nitride w-BN;

further, the boron nitride is hexagonal boron nitride.

5. The iron-based ceramic composite material with an aluminum alloy surface according to claim 1, wherein: the chemical composition of the Fe-based self-fluxing alloy powder is Cr 16.5 wt.%, B0.9 wt.%, si0.8 wt.%, C0.12 wt.%, the balance being Fe and unavoidable impurities.

6. An iron-based ceramic composite coating on the surface of an aluminum alloy is characterized in that: the coating material is the alloying material of claims 1-5; the base material is aluminum alloy;

further, the base material is a 2-series, 4-series, 5-series, 6-series or 7-series aluminum alloy;

furthermore, the base material is 6061 aluminum alloy, and the main phase in the coating comprises Fe₄Al₁₃、Al₃Ti、TiB₂、TiC、TiN、AlN、Cr₂B、Cr₂₃C₆、Al₄C₃。

7. The preparation method of the iron-based ceramic composite coating on the surface of the aluminum alloy is characterized by comprising the following steps of: polishing and cleaning the surface of a base material, grinding Fe-based self-fluxing alloy powder, Ti powder and ceramic powder until the Fe-based self-fluxing alloy powder, the Ti powder and the ceramic powder are uniformly mixed, and presetting a binder on the surface of the base material for laser alloying to obtain a laser alloying coating;

further, the binder is water glass, and the volume ratio of the water glass to the water is 1: 3.

8. The method of preparing a coating according to claim 7, wherein: the laser power is 1-3 kW, the scanning speed is 9-15 mm/s, the laser spot diameter is 3-4 mm, and the lap joint rate is 30-50%;

further, the laser power is 1.7-2.3 kW, the scanning speed is 9-12 mm/s, the spot diameter is 3mm, and the lap joint rate is 30%;

furthermore, the laser power is 2kW, the scanning speed is 12mm/s, the spot diameter is 3mm, and the lap joint rate is 30%.

9. The method of preparing a coating according to claim 7, wherein: in the laser alloying process, argon is used as protective gas, and the gas flow is 8-12L/min;

further, the flow rate of argon gas was 10L/min.

10. The application of the iron-based ceramic composite coating on the surface of the aluminum alloy obtained by the preparation method of any one of claims 7 to 9 in the fields of aerospace, mechanical manufacturing, electronics or transportation.