阅读说明：本技术 一种2014铝合金航空精密轮毂模锻件的无锆毛坯 (No zirconium blank of accurate wheel hub die forging of 2014 aluminum alloy aviation ) 是由 陈丽芳 吴道祥 王正安 林海涛 曾庆华 于 2019-11-28 设计创作，主要内容包括：本发明公开了一种2014铝合金航空精密轮毂模锻件的无锆毛坯,包含0.65％-0.75％Si,Fe≤0.10％,4.00％-4.40％Cu,0.70％-0.90％Mn,0.45％-0.55％Mg,0.04％-0.07％Cr,Zn≤0.20％,Ti≤0.15％。其性能较好,无论是模锻平行晶粒方向以及垂直晶粒方向的屈服强度、抗拉强度和断后伸长率,还是自由锻或者轧环切向以及轴向的屈服强度、抗拉强度和断后伸长率,均高于现有技术中的AMS4133E标准,并且疲劳寿命测试的成绩优异,使用本发明提供的2014铝合金航空精密轮毂模锻件的无锆毛坯,能够极大的提高2014铝合金航空精密轮毂模锻件的产品质量。(The invention discloses a zirconium-free blank of a 2014 aluminum alloy aviation precision hub die forging, which comprises 0.65-0.75% of Si, less than or equal to 0.10% of Fe, 4.00-4.40% of Cu, 0.70-0.90% of Mn, 0.45-0.55% of Mg, 0.04-0.07% of Cr, less than or equal to 0.20% of Zn and less than or equal to 0.15% of Ti. The zirconium-free blank of the 2014 aluminum alloy aviation precision hub die forging piece provided by the invention has the advantages that the performance is good, the yield strength, the tensile strength and the elongation after fracture in the die forging direction parallel to the crystal grains and the direction vertical to the crystal grains, the yield strength, the tensile strength and the elongation after fracture in the free forging or rolling ring tangential direction and the axial direction are all higher than the AMS4133E standard in the prior art, and the performance of a fatigue life test is excellent.)

1. The zirconium-free blank of the 2014 aluminum alloy aviation precision hub die forging is characterized by comprising 0.65-0.75% of Si, less than or equal to 0.10% of Fe, 4.00-4.40% of Cu, 0.70-0.90% of Mn, 0.45-0.55% of Mg, 0.04-0.07% of Cr, less than or equal to 0.20% of Zn and less than or equal to 0.15% of Ti.

2. The zirconium-free blank of the 2014 aluminum alloy aviation precision hub die forging of claim 1, wherein the Si is 0.70%.

3. The zirconium-free blank of the 2014 aluminum alloy aviation precision hub die forging of claim 1, wherein the Fe is 0.08%.

4. The zirconium-free blank of the 2014 aluminum alloy aerospace precision hub die forging of claim 1, wherein the Cu is 4.10%.

5. The zirconium-free blank of the 2014 aluminum alloy aviation precision hub die forging of claim 1, wherein the Mn is 0.80%.

6. The zirconium-free blank of the 2014 aluminum alloy aviation precision hub die forging of claim 1, wherein the Mg is 0.50%.

7. The zirconium-free blank of the 2014 aluminum alloy aviation precision hub die forging of claim 1, wherein the Cr is 0.06%.

8. The zirconium-free blank of the 2014 aluminum alloy aviation precision hub die forging of claim 1, wherein Zn is 0.04%.

9. The zirconium-free blank of the 2014 aluminum alloy aviation precision hub die forging of claim 1, wherein the Ti is 0.02%.

10. The zirconium-free blank of the 2014 aluminum alloy aviation precision hub die forging of claim 1, wherein the Si is 0.72%, the Fe is 0.08%, the Cu is 4.16%, the Mn is 0.80%, the Mg is 0.54%, the Cr is 0.06%, the Zn is 0.04%, and the Ti is 0.03%.

Technical Field

The invention relates to the technical field of manufacturing of aviation precision hub die forgings, in particular to a zirconium-free blank of a 2014 aluminum alloy aviation precision hub die forging.

Background

The large airplane is provided with a typical-specification forge piece which is the largest forge piece in the 2014-high aluminum alloy aviation precision hub die forge piece: and die forging of half wheel (inboard). The half-wheel (inboard) die forging is a precision die forging and is a disc die forging, the maximum outer hub size of a part is phi 593.3 multiplied by 309.1mm, and the maximum outer contour size of the die forging is phi 616.5 multiplied by 314.2 mm.

The parts inside the half wheel cabin are shown in fig. 1 and fig. 2, and fig. 1 is a first side view structure schematic diagram of the 2014 aluminum alloy aviation precision hub die forging provided by the embodiment of the invention; fig. 2 is a schematic side view of a 2014 aluminum alloy aviation precision hub die forging, which is a relatively complex large aluminum alloy forging, the maximum external dimension of the forging is phi 600mm × 310mm, the maximum depth of the cylinder is 240mm, the minimum position of the cylinder wall is only 7.6mm, and the maximum position of the cylinder wall is 16mm, and the forging is a typical deep-cylinder thin-wall part, the basic body of which is a cylinder 12, the upper part of the cylinder 12 is provided with an annular outward extension part 11, an inner concave part is arranged above the outward extension part 11, 9 lugs 14 arranged in an annular manner are arranged at the junction of the inner concave part and the inner wall of the cylinder 12, and the bottom of the cylinder 12 is provided with 9 annular elliptical pits 13, specifically, the part is thin at the bottom of the cylinder, and has 9 uniformly distributed elliptical pits 13 at the same time, and the shape is complex; the upper side of the part is correspondingly provided with 9 lugs 14, and the lugs 14 are high in height, thin in wall thickness, small in inclination and small in vertical projection area, and belong to parts which are difficult to form and easy to have defects.

The half-wheel (inboard) die forging is a precision die forging, namely a disc die forging, and has the advantages of deep cavity, thin wall, high rib, small fillet, more bosses at the inner cavity and the bottom and more complex cavity. The half-wheel (inboard) die forging has a large number of non-machined surfaces, small machining allowance, high surface quality requirement and extremely high dimensional precision requirement; the die forging has deep cavity, high and thin ribs and difficult precision die forging forming; the 2014 alloy is easy to generate coarse grains, and the uniformity of the structure performance is difficult to control; the safety performance requirement of the wheel hub is high, and the comprehensive performance requirement is extremely high. Therefore, the biggest difficulties of the hub die forging are large difficulty in controlling the size and the uniformity of the structure performance.

Therefore, how to provide a zirconium-free blank of the 2014 aluminum alloy aviation precision hub die forging to improve the product quality is a technical problem to be solved by technical personnel in the field.

Disclosure of Invention

In view of the above, the invention aims to provide a zirconium-free blank of a 2014 aluminum alloy aviation precision hub die forging to improve the product quality.

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

a zirconium-free blank of a 2014 aluminum alloy aviation precision hub die forging piece comprises 0.65-0.75% of Si, less than or equal to 0.10% of Fe, 4.00-4.40% of Cu, 0.70-0.90% of Mn, 0.45-0.55% of Mg, 0.04-0.07% of Cr, less than or equal to 0.20% of Zn, and less than or equal to 0.15% of Ti.

Preferably, the Si content is 0.70%.

Preferably, the amount of Fe is 0.08%.

Preferably, the Cu content is 4.10%.

Preferably, the Mn is 0.80%.

Preferably, the amount of Mg is 0.50%.

Preferably, the amount of Cr is 0.06%.

Preferably, the Zn content is 0.04%.

Preferably, the Ti content is 0.02%.

Preferably, the content of Si is 0.72%, the content of Fe is 0.08%, the content of Cu is 4.16%, the content of Mn is 0.80%, the content of Mg is 0.54%, the content of Cr is 0.06%, the content of Zn is 0.04%, and the content of Ti is 0.03%.

The invention provides a zirconium-free blank of a 2014 aluminum alloy aviation precision hub die forging, which comprises 0.65-0.75% of Si, less than or equal to 0.10% of Fe, 4.00-4.40% of Cu, 0.70-0.90% of Mn, 0.45-0.55% of Mg, 0.04-0.07% of Cr, less than or equal to 0.20% of Zn, and less than or equal to 0.15% of Ti. The 2014-aluminum alloy aviation precision hub die forging piece has good performance, the yield strength, the tensile strength and the elongation after fracture of the die forging in the direction parallel to the crystal grains and in the direction vertical to the crystal grains, the yield strength, the tensile strength and the elongation after fracture of free forging, or the yield strength, the tensile strength and the elongation after fracture of a rolling ring in the tangential direction and in the axial direction are all higher than the AMS4133E standard in the prior art, and the fatigue life test performance is excellent.

Drawings

In order to more clearly illustrate the embodiments of the present invention or the technical solutions in the prior art, the drawings used in the description of the embodiments or the prior art will be briefly introduced below, and it is obvious that the drawings in the following description are some embodiments of the present invention, and for those skilled in the art, other drawings can be obtained according to these drawings without creative efforts.

FIG. 1 is a schematic side view structure diagram of a 2014 aluminum alloy aviation precision hub die forging provided in an embodiment of the present invention;

FIG. 2 is a second side view structural schematic diagram of the 2014 aluminum alloy aviation precision hub die forging provided in the embodiment of the invention;

FIG. 3 is a graphical representation of macrosegregation data for alloy # 4;

FIG. 4 is a schematic representation of the as-cast structure of alloy # 1 (normalized median);

FIG. 5 is a schematic representation of the as-cast structure of alloy # 5 (low CuMgSi);

FIG. 6 is a schematic diagram of a three-stage homogenization regime for ingot casting;

FIG. 7 is a schematic view showing coarse phases in the homogenized structure of the 1# alloy;

FIG. 8 is a schematic view showing coarse phases in the homogenized structure of the 5# alloy;

FIG. 9 is a schematic view of a 1# (standard median) diffuse phase scan;

FIG. 10 is a schematic view of a 2# (low Fe + Zr) diffuse phase scan;

FIG. 11 is a schematic drawing of a 3# (low FeMn) diffuse phase scan;

FIG. 12 is a schematic view of an extrusion process route for 5 alloys;

FIG. 13 is a schematic of a solution aging process;

FIG. 14 is a schematic view of a sampling position;

FIG. 15 is a schematic structural view of a high cycle fatigue specimen.

In the above FIGS. 1-15:

the outer extension part 11, the cylinder 12, the oval pit 13 and the lug 14.

Detailed Description

In order to make the objects, technical solutions and advantages of the embodiments of the present invention clearer, the technical solutions in the embodiments of the present invention will be clearly and completely described below with reference to the drawings in the embodiments of the present invention, and it is obvious that the described embodiments are some, but not all, embodiments of the present invention. All other embodiments, which can be derived by a person skilled in the art from the embodiments given herein without making any creative effort, shall fall within the protection scope of the present invention.

Referring to fig. 1 to fig. 15, fig. 1 is a first side view structural schematic diagram of an 2014 aluminum alloy aviation precision hub die forging provided in an embodiment of the present invention; FIG. 2 is a second side view structural schematic diagram of the 2014 aluminum alloy aviation precision hub die forging provided in the embodiment of the invention; FIG. 3 is a graphical representation of macrosegregation data for alloy # 4; FIG. 4 is a schematic representation of the as-cast structure of alloy # 1 (normalized median); FIG. 5 is a schematic representation of the as-cast structure of alloy # 5 (low CuMgSi); FIG. 6 is a schematic diagram of a three-stage homogenization regime for ingot casting; FIG. 7 is a schematic view showing coarse phases in the homogenized structure of the 1# alloy; FIG. 8 is a schematic view showing coarse phases in the homogenized structure of the 5# alloy; FIG. 9 is a schematic view of a 1# (standard median) diffuse phase scan; FIG. 10 is a schematic view of a 2# (low Fe + Zr) diffuse phase scan; FIG. 11 is a schematic drawing of a 3# (low FeMn) diffuse phase scan; FIG. 12 is a schematic view of an extrusion process route for 5 alloys; FIG. 13 is a schematic of a solution aging process; FIG. 14 is a schematic view of a sampling position; FIG. 15 is a schematic structural view of a high cycle fatigue specimen.

The zirconium-free blank of the 2014 aluminum alloy aviation precision hub die forging provided by the embodiment of the invention comprises 0.65-0.75% of Si, less than or equal to 0.10% of Fe, 4.00-4.40% of Cu, 0.70-0.90% of Mn, 0.45-0.55% of Mg, 0.04-0.07% of Cr, less than or equal to 0.20% of Zn and less than or equal to 0.15% of Ti. The 2014-aluminum alloy aviation precision hub die forging provided by the embodiment of the invention has the advantages that the performance is good, the yield strength, the tensile strength and the elongation after fracture of the die forging in the direction parallel to the crystal grains and in the direction vertical to the crystal grains, the yield strength, the tensile strength and the elongation after fracture of the free forging, or the yield strength, the tensile strength and the elongation after fracture of the rolling ring in the tangential direction and in the axial direction are all higher than the AMS4133E standard in the prior art, the fatigue life test performance is excellent, and the product quality of the 2014-aluminum alloy aviation precision hub die forging can be greatly improved by using the zirconium-free blank of the 2014-aluminum alloy aviation precision hub die forging provided by.

Specifically, Si is 0.70%. The Fe content is 0.08%. Cu is 4.10%. Mn is 0.80%. Mg is 0.50%. Cr is 0.06%. Zn is 0.04%. Ti is 0.02%.

Or 0.72% of Si, 0.08% of Fe, 4.16% of Cu, 0.80% of Mn, 0.54% of Mg, 0.06% of Cr, 0.04% of Zn and 0.03% of Ti.

The embodiment of the invention provides a zirconium-free blank of a 2014 aluminum alloy aviation precision hub die forging, which is designed by the following alloy components:

in the AMS4133 standard component range, 5 2014 alloys are designed according to different functions of Cu, Mg, Si, Mn, Fe and Zr in the alloys, and the chemical components of the five 2014 alloys are shown in Table 1, and the Table 1 is shown in the table. The principle of component design: taking the intermediate value of 2014 standard components for the No. 1 alloy; the 2# alloy reduces the Fe content, increases Zr element, and researches the action of the Zr element; the 3# alloy simultaneously reduces Fe and Mn elements and researches the action of the Mn element; the content of Cu, Mg and Si is increased simultaneously in the No. 4 alloy; the 5# alloy simultaneously reduces the contents of Cu, Mg and Si, and the effects of Cu, Mg and Si are compared and analyzed.

The above 5 components were cast, and 1# component was cast into 2 ingots, and the remaining components were cast into 1 ingot each having an ingot size of Φ 178 × 550(37 kg). FIG. 3 is a graphical representation of macrosegregation data for alloy # 4. In the radial direction of the ingot, the content of Cu and Mg elements at the position which is far away from the center 3/4R of the ingot is higher than that at other positions, but the content does not exceed the design components.

TABLE 1 chemical compositions of five 2014 alloys

FIG. 4 is a schematic diagram showing the as-cast structure of alloy # 1 (standard median value), and FIG. 5 is a schematic diagram showing the as-cast structure of alloy # 5 (low CuMgSi). analysis of the as-cast structure of the 5 alloys shows that the ① 1#, 2#, 3#, 4#, and 5# alloys all contain four coarse phases, namely Al2Cu, Al (FeMnSi), Al (CuMgSi), and Mg2 Si. ② 2# (low Fe + Zr) results in a significant reduction in eutectic structure and an increase in bulk Al2Cu phase ③ EDS shows that most of the 2# (low Fe + Zr) alloys contain a higher Zr content, which is higher than 1# (standard median value), and Zr is considered to be mainly dissolved in the aluminum matrix at the as-cast test point.

Determining a soaking system:

and taking the cast structure of 1# standard intermediate value, 4# high CuMgSi and 5# low CuMgSi at a distance of 3/4R from the circle center, and carrying out DSC test. First, the first stage homogenization employed (55 ℃/h ramp) 450 ℃ for 5 h. The second stage homogenization process was initially determined by DSC as: 500 ℃ for 24 h. The third stage homogenization process is initially determined as: 506 ℃ and is temporarily set for 18 h. The 5 kinds of alloy ingots were homogenized by the homogenization process described above, and as shown in fig. 6, fig. 6 is a schematic view of a three-stage homogenization system of the ingots.

The DSC detection was performed on the homogenized ingot to obtain the following conclusions:

(1) homogenizing the alloy ingots of No. 1, No. 2, No. 3 and No. 4 at 500 ℃ for 24h, and homogenizing at 506 ℃ for 18.5h, wherein the peak area is not obviously reduced, which indicates that the third-stage homogenizing effect is not great. Therefore, it is necessary to increase the second homogenization temperature to make the soluble phase redissolved as much as possible during the second homogenization.

(2) After homogenizing the 5# alloy cast ingot at 500 ℃ for 24h +506 ℃ for 18.5h, most of the Al2Cu phase is redissolved; the onset of the residual phase redissolution endotherm is shifted to 520 ℃ with a peak 527 ℃ close to the melting point of the Q phase. Therefore, the homogenization degree basically meets the 5# alloy.

FIG. 7 shows the coarse phase in the homogenized structure of the 1# alloy. FIG. 8 shows the coarse phase in the homogenized structure of the 5# alloy. After the 5 alloys are homogenized, the Al2Cu phase is not completely redissolved, and the non-redissolved Al2Cu and the insoluble Al (FeMnSi) are still retained at the grain boundary. Although DSC of alloy # 5 showed that the Al2Cu peak had been eliminated, a small portion of bulk Al2Cu (about 5 microns in size) remained in the structure.

The type and number density analysis of the first-stage homogenization precipitated dispersoid phase is respectively carried out on the 1# (standard intermediate value), the 2# (low Fe + Zr) and the 3# (low FeMn) alloys. Table 2 shows the experimental protocol for dispersed phase analysis.

TABLE 2 Dispersion phase analysis protocol

Selection of alloys	Status of state	Purpose(s) to
			1# (Standard median)	Water quenching at 450 deg.c/5 hr +505 deg.c/24 hr	Analysis of dispersed phase species, number density
2# (Low Fe + Zr)	Water quenching at 450 deg.c/5 hr +505 deg.c/24 hr	Analysis of dispersed phase species, number density
			3# (Low FeMn)	Water quenching at 450 deg.c/5 hr +505 deg.c/24 hr	Analysis of dispersed phase species, number density

Referring to fig. 9, 10 and 11, the dispersoid phase scan photographs of 1# (standard median), 2# (low Fe + Zr) and 3# (low FeMn) are shown in sequence. The types of the dispersoids in the alloys # 1, # 2 and # 3 were identified and the number density was calculated from the scanning photographs of the dispersoids, and the results are shown in Table 3. Table 3 shows the results of laboratory homogenization analysis of dispersed phases for Nos. 1#, 2#, and 3 #. Table 3 shows the results of laboratory homogenization analysis of dispersed phases for Nos. 1#, 2#, and 3 #.

Table 31 #, 2#, 3# laboratory homogenized precipitation dispersion phase analysis results

The homogeneous precipitated dispersoids were analyzed and the following conclusions were obtained:

(1) the Mn-containing dispersoids precipitated in the alloys No. 1 and No. 3 are the same in type, and are Al (Mn2Si) and Al (Mn). Little Al-Zr dispersed phase was observed alone in alloy # 2, since first stage homogenization 450/5h was not the optimum precipitation temperature for Al3 Zr.

(2) In the process of heating the 2# alloy from room temperature to 450 ℃, Zr may be precipitated firstly; in the heat preservation stage at 450 ℃, Al3Zr can become a nucleation point for precipitation of an Al-Mn-containing phase, so that the density and the size of the Al-Mn-containing phase are larger and smaller.

(3) The Mn content in No. 3 is low, the dispersoid phase density of Al (Mn2Si) and Al (Mn) is obviously lower than that of No. 1 and No. 2, and the size is larger.

And (3) carrying out test verification:

the alloy performance can be better compared by adopting extrusion deformation. The 5 alloys are subjected to hot extrusion, and the existing die (the section size is 25mm multiplied by 125mm) is selected to ensure that the size of the extruded plate strip can meet the tensile and fatigue performance tests. The extrusion process route is shown in fig. 12, and after the final extrusion, the head section is cut off by 0.4m, the tail section is cut off by 1m, and the rest about 2m can be used.

According to the AMS 2772 standard recommended system, the solution temperature is 502 ℃ multiplied by the heat preservation time for 3.5 h; the aging process is 177 ℃ multiplied by 8 h. Cutting an extruded plate with the length of 400mm, and carrying out solid solution aging, wherein the aging treatment of the solid solution treatment is heating along with a furnace. The solution aging process is shown in fig. 13, and the sampling position is shown in fig. 14.

And selecting the L-S surface of the extruded plate strip to perform EBSD structure analysis, comparing with the extruded EBSD, and analyzing the influence of the solid solution heat treatment on the structure. All 5 alloys showed elongated ribbon-like structures and the degree of recrystallization was not evident, but slightly different. The detailed description is as follows: after the alloys 1#, 3#, 4# are subjected to solid solution for 3 hours and aged, fine recrystallized grains in an extruded state do not grow obviously after the solid solution for 3 hours. After the 2# and 5# alloys are subjected to solid solution and aging for 5 hours, fine recrystallization occurs and slight growth occurs. Comprehensively considered, the strength of 5h solid solution and aging is higher, and the plasticity is better; however, an increase in the solution time leads to an increase in the fraction of recrystallized grains, which may be detrimental to the fatigue life.

Table 4 lists the recrystallization fraction of the solid solution age state of the 5 alloys counted by the EBSD recrystallization module. The EBSD picture of the LS surface of the sample contains a small number of crystal grains, and the EBSD picture of the TS surface of the sample contains a large number of crystal grains; therefore, the recrystallization fraction of the TS surface is statistically accurate. Table 4 shows the recrystallization fractions of the T6 states of the 5 alloys.

TABLE 45 recrystallization fractions of alloys in the T6 state

The residual phases of the 5 alloys were analyzed:

taking alloy # 1 as an example, different phase types are distinguished by contrast differences: bright white is Al2Cu, light gray is Al (femnsi), dark gray is Al (cumgsi). Table 5 lists the statistics of the three residual phases. Table 5 compares the content of all residual phases.

TABLE 5 comparison of all residual phase contents

(1) First, Al2Cu was analyzed:

the Cu content in 1# (4.49Cu) and 3# (4.48Cu) is medium, and the Cu content in 5# (4.2Cu) alloy is low, which is consistent with the area percentage rule of Al2Cu in the alloy.

2# (4.3Cu) has a low Cu content, but a high residual Al2Cu content. In the as-cast structure of the 2# alloy, the eutectic structure was less, and the bulk Al2Cu was more. Therefore, the Zr is added into the 2# alloy, so that the shape of the Al2Cu tends to be blocky; in subsequent homogenization, the dissolution back of bulk Al2Cu is slower than in the eutectic structure, thus more Al2Cu phase remains in the final T6 state.

The highest Cu content in # 4 (4,75Cu) and the homogenized DSC and phase statistics above indicated that the Al2Cu phase content was also higher.

(2) Next, al (femnsi):

the Al (FeMnSi) phase content is determined by the Mn content first, e.g., 3# (0.12Fe, 0.62Mn) with the least Al (FeMnSi) phase content.

The Al (FeMnSi) phase content is determined by the Fe content, and the Mn content is the same as that of the four phases of 1# (0.19Fe, 0.8Mn), 2# (0.15Fe, 0.8Mn) and 5# (0.09Fe, 0.8Mn), and the Al (FeMnSi) content is in positive correlation with the Fe content basically.

The FeMn content in No. 4 (0.09Fe, 0.77Mn) is lower than that in No. 5 (0.09Fe, 0.8Mn), and the Al (FeMnSi) phase content is correspondingly lower than that in No. 5.

(3) Finally, al (cumgsi):

since the Al (cumgsi) phase has no significant contrast in SEM pictures, the statistical data confidence is low compared to Al2Cu and Al (femnsi), but 5 alloys can still be compared with each other.

The content of CuMg in No. 4 (4.75Cu, 0.8Mg) is the highest, and the corresponding Al (CuMgSi) phase has the most residue. The content of CuMg in # 5 (4.2Cu, 0.68Mg) was the lowest, and no Al (CuMgSi) phase was found. It is considered that the Cu and Mg contents are substantially in a positive correlation with the Al (CuMgSi) phase content.

(4) Analysis of all residual phases:

the percentage of residual phase area for the 5 alloys was analyzed to conclude that: the content of Al2Cu residual phase is reduced by reducing the Cu content; reducing the content of Al (FeMnSi) residual phase by reducing the content of Fe and Mn; the content of Al (CuMgSi) residual phase is reduced by reducing the content of Cu and Mg.

Table 6 shows the room temperature tensile properties data of the extrudate T6 (3h solid solution + aging) and Table 7 shows the room temperature tensile properties data of the extrudate T6(5 h solid solution + aging), wherein △ σ 0.2(L-T), △ σ b (L-T) and △ δ (L-T) show the difference in yield strength, the difference in tensile strength and the difference in elongation after fracture in the extrusion direction and in the transverse direction, respectively, and Table 8 shows the influence of the aging time of 3h and 5h solid solution on the properties.

TABLE 6 tensile properties at room temperature in the T6 state (502 ℃ C.. times.3 h +177 ℃ C.. times.8 h)

TABLE 7 tensile properties at room temperature in the T6(502 ℃ C.. times.5 h +177 ℃ C.. times.8 h) state

TABLE 8 Effect of aging times 3h and 5h on Properties

By room temperature tensile property analysis, the following conclusions were obtained:

(1) the T6 state performance of the 5 alloys is higher than AMS4133E standard. The solid solution for 5h improves the strength by 5-20MPa and the elongation by 1-4.5% compared with the solid solution for 3h, probably because the long-time solid solution increases the re-dissolution degree of the residual Al2Cu phase and enhances the aging strengthening.

(2) The 3# alloy has low Mn content, the recrystallization inhibiting effect is weakened, the cubic texture is enhanced through solution treatment, the strength in the extrusion direction is weaker than that of other alloys, and the performance difference between the extrusion direction and the transverse direction is small.

The high cycle fatigue life test is carried out on the extrusion plate belt, and the stress level is 300MPa, the stress ratio is-1, and the frequency is 50 Hz. Each of the L-direction and T-direction 3 parallel specimens of each alloy was an unnotched smooth specimen, and the dimensions of the high cycle fatigue specimens used in this experiment are shown in fig. 15. Table 9 shows the T6 temper (502 ℃ C.. times.5 h +177 ℃ C.. times.8 h) fatigue properties of 5 alloys.

TABLE 95 alloy T6(502 ℃ C.. times.5 h +177 ℃ C.. times.8 h) fatigue properties

Fatigue life in the extrusion direction (L direction) (60-130 ten thousand times) is generally higher than in the transverse direction (T direction) (10-90 ten thousand times). The 5# (low CuMgSi) alloy has the highest life in the extrusion direction, but the lowest life in the transverse direction. The 4# (high CuMg) alloy had the least difference in life in the extrusion direction and transverse direction.

The fatigue fracture is analyzed, the crack propagation areas of the 5 alloys are not very different, and the initiation areas are very different. Therefore, the crack source is observed with emphasis: (1) fracture macro morphology (2), crack initiation position (3), and phase composition of crack initiation position.

Fatigue fracture of 2# L-2 specimen (76 ten thousand cycles): mg, O enriched regions were observed and no other coarse macrophases were found.

Fatigue fracture of 2# T-1 specimen (53 ten thousand cycles): mg, O enriched regions were observed and no other coarse macrophases were found.

For the 2# T-1 sample, cracks clearly initiated in the coarse phase particles. The composition shows that the coarse phase is rich in Mg and O, contains Si, and other elements are uniformly distributed. Around the coarse phase, a Cu-rich coarse phase exists.

In a fatigue fracture of a 2# L-2 sample (99 ten thousand weeks), Mg and O elements are enriched at a crack source; meanwhile, a small amount of Fe, Mn and Si enrichment points exist.

The fatigue life analysis is more complex, and relates to the type, quantity, size, distribution, recrystallization and mechanical properties of a matrix of the coarse phase. In addition, fatigue data is scattered and analysis is also difficult. From the experimental data, the fatigue life analysis for the 5 alloys is as follows:

initiation of fatigue cracks: according to fracture analysis of 2# (containing Zr) and 5# (low CuMg) alloys, crack sources are generated at Mg-rich and O-rich positions, and Fe, Mn and Si enrichment points exist around the crack sources. Indicating that crack initiation is more sensitive at MgO than at the coarse phase. The batch of alloy is cast in Su's institute, and the conditions of degassing and impurity removing equipment are limited, which may cause excessive MgO in the alloy.

Although cracks were found to be initiated generally at MgO, comparing 4# (high CuMg) and 5# (low CuMg) alloys found: no. 5 (low CuMg) extrusion direction fatigue life is the highest, but the transverse life is not obviously improved. The 4# (high CuMg) transverse fatigue performance is highest, but the extrusion direction life is not obviously reduced. It was preliminarily believed that MgO was consistent in fatigue crack initiation time, and the effect of the coarse phase may be reflected in crack propagation.

Through comprehensive analysis, a component optimization scheme is provided:

(1) a Zr-free alloy: 4.1% Cu, 0.4% Mg, 0.1% Fe, 0.7% Si, 0.8% Mn. When the controllable deviation of the furnace temperature is +/-5 ℃, the homogenization process adopts: 450 ℃ X5 h +505 ℃ X30 h.

(2) Adding Zr alloy: 4.1% Cu, 0.4% Mg, 0.1% Fe, 0.7% Si, 0.6% Mn, 0.1% Zr. When the controllable deviation of the furnace temperature is +/-5 ℃, the homogenization process adopts: 400 ℃ X10 h +450 ℃ X5 h +505 ℃ X30 h.

As shown in Table 10, Table 10 is a control chart of the composition of the cast product in actual production.

TABLE 102014 ingot casting Final composition

The previous description of the disclosed embodiments is provided to enable any person skilled in the art to make or use the present invention. Various modifications to these embodiments will be readily apparent to those skilled in the art, and the generic principles defined herein may be applied to other embodiments without departing from the spirit or scope of the invention. Thus, the present invention is not intended to be limited to the embodiments shown herein but is to be accorded the widest scope consistent with the principles and novel features disclosed herein.