IRON-BASED ALLOY COMPOSITION, PARTS PRODUCED FROM THIS COMPOSITION AND PRODUCTION METHOD

Abstract

An iron-based alloy composition includes 0.28-0.34% C, max 0.25% Si, max 0.8% Mn, 0.85-0.95% Cr, 1.10-1.50% Ni, 0.41-0.50% Mo, 0.001-0.007% B, 0.002-0.03% Nb, and balanced amount of Fe and inevitable impurities. Moreover, parts with a hardness of at least 480 HB, a tensile strength of at least 1700 MPa, a total elongation of at least 7% and an impact strength of at least 16 J are obtained by the iron-based alloy composition and a production method of the parts.

Claims

1. An iron-based alloy composition developed for producing a hot formed armor steel, comprising 0.28-0.34% carbon, max 0.25% silicon, max 0.8% manganese, 0.85-0.95% chromium, 1.10-1.50% nickel, 0.41-0.50% molybdenum, 0.001-0.007% boron, 0.002-0.03% niobium by weight and a balanced amount of iron and inevitable impurities.

2. The iron-based alloy composition according to claim 1, comprising one or more elements selected from the group containing trace amounts of phosphorus, sulfur, copper, aluminum, tungsten, cobalt, titanium, oxygen, hydrogen, and nitrogen.

3. A hot formed armor steel, comprising 0.28-0.34% carbon, max 0.25% silicon, max 0.8% manganese, 0.85-0.95% chromium, 1.10-1.50% nickel, 0.41-0.50% molybdenum, 0.001-0.007% boron, 0.002-0.03% niobium by weight and a balanced amount of iron and inevitable impurities in a composition of the hot formed steel.

4. The hot formed armor steel according to claim 3, comprising one or more elements selected from the group containing trace amounts of phosphorus, sulfur, copper, aluminum, tungsten, cobalt, titanium, oxygen, hydrogen, and nitrogen.

5. The hot formed armor steel according to claim 3, wherein the hot formed armor steel has a hardness of at least 480 HB, a tensile strength of at least 1700 MPa, a total elongation of at least 7%, and/or an impact strength of at least 16 J.

6. The hot formed armor steel according to claim 3, comprising at least 90% martensite in a microstructure of the hot formed armor steel.

7. A hot formed armor steel production method, comprising the following steps: i. an ingot or slab casting of an alloy comprising 0.28-0.34% carbon, max 0.25% silicon, max 0.8% manganese, 0.85-0.95% chromium, 1.10-1.50% nickel, 0.41-0.50% molybdenum, 0.001-0.007% boron, 0.002-0.03% niobium by weight, and balanced amounts of iron and inevitable impurities to obtain an ingot or a slab, ii. hot rolling the slab or the ingot into a plate, iii. plate cooling and cutting, iv. applying a primary heat treatment to cut the plate, v. forming the plate by pressing in a cooled tool, vi. applying a secondary heat treatment to formed steel parts.

8. The hot formed armor steel production method according to claim 7, comprising one or more elements selected from the group containing trace amounts of phosphorus, sulfur, copper, aluminum, tungsten, cobalt, titanium, oxygen, hydrogen, and nitrogen in the step i.

9. The hot formed armor steel production method according to claim 7, comprising heating the slab or the ingot to above 1050° C. for at least 4 hours.

10. The hot formed armor steel production method according to claim 7, comprising cooling the plate down to 2° C./s or slower, a microstructure of ferrite+perlite, bainite, or a mixture of phases and obtaining a plate with a hardness scale below 300 HB, a heating process of the plate above 300° C. and transforming a microstructure of the plate into a tempered martensite, provided that the heating process is performed faster without cooling in the step iii.

11. The hot formed armor steel production method according to claim 7, wherein the primary heat treatment for cutting the plate is performed by heating the plate to a temperature below 1000° C. and above AC3 for at least 10 minutes in the step iv.

12. The hot formed armor steel production method according to claim 7, comprising forming the plate by cooling the plate to a temperature of 300° C. or less at a rate of over 4° C./s in the step v.

13. The hot formed armor steel production method according to claim 7, comprising tempering of the formed steel parts in the step vi by applying the secondary heat treatment at a temperature of 250° C. or less, and obtaining at least 90% martensitic microstructure.

14. The hot formed armor steel production method according to claim 7, comprising tempering of the formed steel parts at a temperature between 140° C.-200° C. by applying the secondary heat treatment for 2-8 hours and obtaining at least 90% martensitic microstructure in the step vi.

15. The hot formed armor steel production method according to claim 7, wherein three-dimensional steel parts obtained after the step vi has a hardness of at least 480 HB, a tensile strength of at least 1700 MPa, a total elongation of at least 7%, and/or an impact strength of at least 16 J.

16. The hot formed armor steel production method according to claim 7, comprising cleaning surfaces of the formed steel parts before or after the step vi.

Description

BRIEF DESCRIPTION OF THE DRAWINGS

[0057] FIGS. 1A-1F The phase transformation graphs obtained in the method according to the invention during cooling process at the cooling rates of the C-001 material produced as hot rolled plate (Ms: Martensite starting temperature, Mf: Martensite final temperature, Bs: Bainite starting temperature)

[0058] FIG. 2 Electron microscope image of the C-001 sample after hot forming and tempering processes

DETAILED DESCRIPTION OF THE EMBODIMENTS

[0059] Alloy composition and production method for the production of hot-formed armor steel according to the invention is explained in this detailed description in order to better understand the subject with its preferred embodiments, and in a way that does not have any restrictive effect.

[0060] The invention is predicated on the development of an iron-based alloy composition for obtaining three-dimensional parts from armor steel hard enough to endure hot forming and which has a ballistic protection feature, and the optimization of the hot forming method using this composition.

[0061] Iron-based alloy composition according to the invention basically includes;

[0062] 0.28-0.34% carbon (C),

[0063] max 0.25% silicon (Si),

[0064] max 0.8% manganese (Mn),

[0065] 0.85-0.95% chromium (Cr),

[0066] 1.10-1.50% nickel (Ni),

[0067] 0.41-0.50% molibden (Mo),

[0068] 0.001-0.007% boron (B),

[0069] 0.002-0.03% niobium (Nb), and

[0070] balanced amounts of iron and inevitable impurities.

[0071] According to one embodiment of the invention, the iron-based alloy composition also may contain one or more unenviable elements selected from the group containing phosphorus (P), sulfur (S), copper (Cu), aluminum (Al), tungsten (W), cobalt (Co), titanium (Ti), oxygen (O), hydrogen (H), nitrogen (N).

[0072] Within the scope of the invention, the alloy formed with the chemical composition described above is transformed from liquid steel form to solid steel form by ingot casting or continuous casting, thus, casting process is performed and the steel is casted into ingot or slab in three-dimensional armor steel hot forming method. The slab or ingot is heated above 1050° C.—preferably to 1200° C.—for at least 4 hours and then hot rolled into a plate.

[0073] Hot rolled plate is cooled down to 2° C./s or slower; therefore, its microstructure includes ferrite+pearlite, bainite, or a mixture of these phases. The plate is heated above 300° C. and its microstructure is transformed into tempered martensite, provided that the process is performed faster without cooling process. In conclusion, if the desired slow cooling rates are achieved, a plate with a hardness scale of 300 HB is obtained.

[0074] Hot rolled plate is cut into desired forms by means of CNC, flame, water jet, laser, saw, etc., and the cut plates are subjected to primary heat treatment for at least 10 minutes by heating to a temperature below 1000° C., and above Ac.sub.3. The heated plate, thereafter, is placed in a water-cooled tool in a press while it is still hot.

[0075] The hot plate is shaped by cooling it to a temperature of 300° C. or below so that a martensitic microstructure is obtained at a speed above 4° C./s by means of the force applied by the press and the water-cooled tool in the press.

[0076] Three-dimensionally formed steel part is removed from the tool and tempered by applying a secondary heat treatment at a temperature of 250° C. or below, and tempered martensitic microstructure is obtained. The part surface is cleaned by sandblasting, polishing, etc.

[0077] Three-dimensional steel parts produced by means of the method described above provide hardness of at least 480 HB, tensile strength of at least 1700 MPa, total elongation of at least 7% and/or impact strength of at least 16 J, and can be used as armored parts with ballistic resistance.

[0078] By means of the recommended method according to the invention, the plate in the chemical composition developed for the steel alloy is produced with a microstructure consisting of ferrite+perlite, bainite or a mixture of these phases, after a cooling process of 2° C./s or slower upon the hot rolling. The microstructure of plate is transformed into tempered martensite when performed faster by heating the plate up to a temperature of 300° C. and above, without the need for a cooling process. The plate produced in this way has a hardness scale below 300 HB and does not yet have a ballistic resistance, making it easier to be cut in the desired form. The microstructure of the plates cut in desired sizes is transformed into martensite by means of being heated to a temperature below 1000° C. and austenitized and then placed in the tool for the forming process, and obtaining a three-dimensional form in the tool with the help of a press and cooling the tool from outside with water. The austenite begins to transform into martensite at a temperature slightly above 300° C. The part with the desired form can be removed from the tool below this temperature. The surface of the three-dimensional steel part can be flattened via sanding and polishing, after a cooling process at room temperature. Surface cleaning process refers to the cleaning of the surface up to a depth of 100 microns. The part is re-heated and tempered by applying heat treatment at a temperature below 250° C. for at least 1 hour at the final phase of the method. Surface cleaning can be applied after tempering as well.

[0079] In the recommended method according to the invention, the chemical composition of the developed steel has been designed in such a way that its microstructure can transform into martensite at cooling rates of 4° C./s or higher during cooling, and thus martensitic structure can be obtained at relatively low cooling rates observed in thick-sectioned parts. Produced steel includes at least 90% martensite microstructurally after hot forming and press hardening processes. The carbon amount in the designed iron-based alloy composition is between 0.28%-0.34%, providing a high weldability in the produced steel. Its manganese amount is below 0.8% in order to prevent segregation. The chromium amount is limited between 0.85%-0.95% so as to delay the perlite formation during cooling process, and to ensure high hardenability. The nickel amount is optimized between 1.10%-1.50% and the molybdenum content between 0.41% and 0.50% in order to increase the hardening depth and hardenability features to provide ballistic properties. Silicon content is limited below 0.25% in order to prevent the formation of silicon oxide in hot rolling and heat treatment processes.

[0080] Welding the armor steels including high amount of carbon with other parts in the vehicle production causes defects in the weld zone. For this reason, this process is not desired by armored vehicle manufacturers. Furthermore, cracks may occur due to thermal stress that occurs during tool cooling in high carbon steels and stresses caused by Bain strains occurring with martensite transformation. The present invention provides the necessary armor feature with a carbon amount of 0.28%-0.34% to eliminate these problems.

[0081] It is known that steels with high manganese content convert into an inhomogeneous structure due to manganese segregation after hot rolling. Even though manganese intensifies the hardenability of the steel, crack formation in continuous casting during steel production causes hardships such as routing, etc. For this reason, manganese amount was limited below 0.8% in the present invention. Furthermore, high silicon content causes silicon-based oxide formation on the surface of the steel after hot rolling or during heat treatment. This oxidation cannot be removed by acid and sanding. It, therefore, constitutes a problem for later coating or painting. Fayalite (Fe.sub.2SO.sub.4) is formed during hot rolling of steels containing silicon, and bonds with FeO. Since this is a strong bond, it makes it difficult to remove the oxide and causes the formation of red scale. The scale cannot be removed by traditional oxide removal methods, causing defects in these areas and decreasing the paint coating workability. It is known in the present art that there is no Fe.sub.2SO.sub.4 compound formation, which causes red scale formation when below 0.25%. Hard oxide layers formed on the surface during hot forming damage and reduce the lifetime of the tool, besides the paint coating workability. Therefore, the silicon amount of the alloy developed in compliance with hot forming was kept below 0.25%. Hereby, its surface properties are also improved. However, silicon affects solid solution hardening positively and raises hardenability. Silicon can be used to prevent carbide formation in steels. Decreasing of hardenability dramatically affects the ballistic properties by enabling formation of unwanted phases during cooling. Hence, 0.41% molybdenum alloying is also employed to increase hardening. Hence, the alloy according to the invention is unique on this sense.

[0082] Iron-based alloy composition developed according to the invention has a structure that can provide ballistic properties with hot forming and tool cooling, and subsequent tempering. Due to low amount of silicon, no red scale formation is observed on the surface of the steel, and the oxidation layer can be, thereby, easily removed. Therefore, it is a material with high paint coating workability. The hardness scale of the final product is at least 480 HB, its tensile strength is typically at least 1700 MPa, its total elongation value is at least 7%, and the notched impact toughness value is at least 16 J at the room temperature.

[0083] Tests and analyzes were carried out with hot formed armor steel part samples obtained within the scope of the invention, and comparative results were recorded and presented in tables below.

[0084] Compositions of the hot formed armor steel samples developed within the scope of the invention are presented in Table 1. Phase transformation at different cooling rates for C-001 alloy is demonstrated in FIGS. 1A-1F. It is evident that bainite is formed prior to martensite transformation when the alloy is cooled down to 2° C./s or slower. Therefore, before hot forming process, the material should be cooled down to 2° C./s or slower after hot rolling in order to be easily cut to the desired dimensions. The hardness scale of the material produced by cooling in this manner is below 300 HB. Mechanical properties of armor steels produced by heating at 900° C. for 10 minutes and via cooling in the tool and subsequent tempering are presented in Table 2. It is evident that the hardness scale above 500 HV are obtained after hot forming and tempering processes. Ballistic performance values of different alloys produced by hot forming and tempering are given in Table 3 and Table 4 after being tested with different ammunitions. The scanning electron microscope image of the C-001 sample after hot forming and tempering processes is given in FIG. 2. It is evident that a martensitic structure has been obtained. It was observed that the ballistic performance of the H009 alloy—whose composition is presented in Table 1—is not sufficient, even though it has the highest impact toughness value. C and Mo contents of the H009 alloy are lower than the other alloys, while its Mn content is slightly higher. The H010 alloy, which is very similar to the H009 alloy and has only a slightly higher C amount, has shown a high ballistic performance against the 7.62×51 Nato Ball ammunition, though having a lower thickness when compared to H009 alloy. Therefore, H009 alloy is excluded from the patent scope. H010 alloy with a slightly higher carbon amount did not display the desired performance on the ballistic tests performed with the 5.56×45 mm SS109 ammunition. Hence, the C-001 alloy is developed by increasing the Cr, Mo, B and Nb amounts of this alloy and decreasing the Mn amount albeit. This alloy is produced via vacuum melting method, differing from other alloys being produced by melting under Ar protection under atmospheric conditions. Therefore, it is ensured that amounts of N, O, H elements and the relevant inclusions are reduced in steel, and the casting cavities are largely eliminated. The ballistic performance of the C-001 alloy produced in different thicknesses are demonstrated in the Table 4 after ballistic tests performed with 7.62×51 Nato Ball ammunition. It is evident that this alloy provides ballistic strength even at lower thicknesses compared to other alloys. The ballistic performance of the C-001 alloy is demonstrated in the Table 5 after ballistic tests performed with 5.56×45 SS109 ammunition, and the desired protection level is reached. For this reason, C-001 alloy is a patented chemical composition. Nonetheless, H-009.5 alloy is also patented as it provides the desired protection level in 7 mm thickness. Due to similar production methods, it is regarded that the N, O, H amounts of the H009.5 alloy are similar to the H010 alloy. C-001 alloy is produced by vacuum melting method. Therefore, the N, O, H amounts are lower. The restricted amounts of N, O, H elements are, thus, considered to be ineffective regarding the ballistic performance. However, due to phenomena such as hydrogen embrittlement and grain boundary corrosion that may arise problems over time, it is aimed to keep the amount of these elements low even if no effect on ballistic performance is observed. No significant effect of Al, S, P elements have been observed against ballistic performance. In the steels produced at the present time, boron addition is limited to 10-20 ppm (Sharma, M., Ortlepp, I., & Bleck, W. (2019). Boron in Heat-Treatable Steels: A Review steel research international, 90(11), 1900133). Nevertheless, it is possible to achieve a higher amount of boron addition with necessary precautions. There are different reasons for this limitation. Considering the manufacturability, in the boron addition phase of the steel, nitrogen forming elements should be bound with nitrogen and B.sub.2O.sub.3 formation should be prevented by keeping the amount of BN and oxygen low. This problem was tried to be avoided with the addition of high amount of Nb (246 ppm) and Al (160 ppm) in C-001 steel and limiting the N (59 ppm) and O (45 ppm) amounts. Otherwise, it is inevitable for free B atoms to form BN. Boron carbides can be formed in boron added steels to some extend even when deemed protection is applied, but their stability is low and they dissolve at temperatures above 800° C. Excessive boron addition (>80 ppm) causes hot shortness. It is possible to work with lower amounts of boron in terms of manufacturability. Furthermore, boron element can be diverged from steel in the phenomenon called “boron fade” when being heated above 900° C. The hardenability is also affected in such a case. This patent relates to thick-sectioned parts; thus, this risk may occur at near-surface, just as decarbonization. According to thermodynamic calculations, up to 41.9 ppm B can dissolve in austenite. This amount may rise up to 97.4 ppm in delta ferrite during solidification. Interstitially or substitutionally dissolving boron generally segregates to grain boundaries and near regions. It increases the hardenability by delaying ferrite or perlite nucleation in these regions. Regarding the hardenability, the boron being in dissolved form or in fine precipitates is considered suitable for a high hardenability. Hardenability decreases with excessive boron and coarse boron-based carbide formation. Nevertheless; in the present case, boron carbides are regarded as dissolved due to quenching process performed in approximately 1000° C. There are publications in terms of toughness regarding that the boron element at the grain boundary decreases the toughness, while there are also papers claiming that it has no effect at all Toughness is generally dependent on steel alloy and is highly related to the toughness level expected from steel. For example, if there is no Al, Ti, Nb and similar elements that can form nitride, BN precipitates can cause austenite grain coarsening and reduce toughness. Therefore, the minimum-maximum amounts of these elements have been determined based on the alloys demonstrated in the Table 1. W, Co, Cu, Ti, Al, S, P elements were observed in trace amounts in 4 different alloys, and no effect on ballistic resistance have been observed. The relevant values per alloy are also given in Table 1.

TABLE-US-00001 TABLE 1 Hot formed armor steel compositions measured by OES before hot stamping process Amount of Element Sample (by weight - %) Code C Si Mn P S Cr Mo Ni W Co Cu H009 0.26 0.11 0.88 0.080 0.020 0.79 0.39 1.27 0.025 0.01 0.045 H009.5 0.30 0.22 0.73 0.080 0.020 0.88 0.41 1.38 0.034 <0.005 0.083 H010 0.34 0.19 0.78 0.070 0.020 0.77 0.37 1.25 0.018 <0.005 0.08 C-001 0.33 0.18 0.70 0.007 0.001 0.91 0.43 1.25 0.029 <0.005 0.064 Amount of Element Sample (by weight - %) Code Ti Al B Nb N O H Fe H009 0.003 0.0045 0.0013 0.0141 n/a n/a n/a Rest H009.5 <0.002 0.0060 0.0023 0.0095 n/a n/a n/a Rest H010 0.003 0.0038 0.0012 0.0156 0.0114 0.0144 0.000058 Rest C-001 <0.002 0.0160 0.0064 0.0246 0.0059 0.0045 0.000088 Rest

TABLE-US-00002 TABLE 2 Mechanical properties of developed armor steels after hot forming and tempering. Impact Tensile Total Micro Toughness Strength Elongation Hardness [J] [MPa] [%] [HV] Sample Std Std Std Std Code* Mean Deviation Mean Deviation Mean Deviation Mean Deviation H009 64 14 1606 4.5 10.9 1.2 512 14 H010 41.1 0.70 2074 23.1 11.30 1.2 584 21 C-001 41.8 0.2 1838.6 37.1 11.1 3.0 553 11

TABLE-US-00003 TABLE 3 Ballistic test results of the H009, H009.5 and H010 materials produced by hot forming and tempering. Measured Measured 7.62 mm × 51 5.56 mm × 45 Sample Sample Nato Ball Nato SS109 Sample Size Thickness Speed Speed Penetration Firing Code [mm] [mm] [m/s] [m/s] after Test 1 H010 100 × 200 6.0 845.74 — NONE 2 — 912.82 NONE 3 — 959.76 OBSERVED 1 H009.5 100 × 200 7.0 842.71 — NONE 2 — 904.53 NONE 3 — 956.99 NONE 1 H009 100 × 195 6.4 839.72 — OBSERVED 2 857.02 — OBSERVED 3 847.08 — OBSERVED

TABLE-US-00004 TABLE 4 Ballistic test results of C-001 material produced by hot forming and tempering in different thickness. Sample Thickness Rate of 7.62 mm × 51 Nato Ball code [mm] fire Speed [m/s] Test result C-001 5.7 1 840 No penetration 2 843 No penetration C-001 6.0 1 830 No penetration

TABLE-US-00005 TABLE 5 Ballistic test results of the C-001 material produced by hot forming and tempering in 6.5 mm thickness. Measured Measured 7.62 mm × 51 5.56 mm × 45 Sample Sample Nato Ball Nato SS109 Sample Size Thickness Speed Speed Penetration Firing Code [mm] [mm] [m/s] [m/s] after Test 1 C-001 100 × 200 6.5 965 NONE 2 — 956 NONE 3 — 955 NONE 4 962 NONE 5 963 NONE 6 971 OBSERVED 7 842 NONE