JP6021094B2 - High-strength non-heat treated steel material excellent in strength, ductility and toughness and method for producing the same - Google Patents
High-strength non-heat treated steel material excellent in strength, ductility and toughness and method for producing the same Download PDFInfo
- Publication number
- JP6021094B2 JP6021094B2 JP2010249102A JP2010249102A JP6021094B2 JP 6021094 B2 JP6021094 B2 JP 6021094B2 JP 2010249102 A JP2010249102 A JP 2010249102A JP 2010249102 A JP2010249102 A JP 2010249102A JP 6021094 B2 JP6021094 B2 JP 6021094B2
- Authority
- JP
- Japan
- Prior art keywords
- steel
- strength
- ductility
- steel material
- toughness
- Prior art date
- Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
- Expired - Fee Related
Links
Landscapes
- Heat Treatment Of Steel (AREA)
Description
本発明は、建造物や橋梁等の構造物、自動車の足回り鋼材、機械用歯車等部品に使用される鋼材であって、特に高強度且つ高延性で、衝撃エネルギー吸収能に優れた厚鋼板や棒鋼・鋼線等の非調質鋼材およびその製造方法に関するものである。
The present invention is a steel material used for parts such as structures such as buildings and bridges, undercarriage steel materials for automobiles, mechanical gears, etc., and is particularly a steel plate having high strength and high ductility and excellent impact energy absorption capability. The present invention relates to non-heat treated steel materials such as steel bars and steel wires, and methods for producing the same.
近年、構造物の大型化や自動車部品の軽量化に伴って、これまで以上に強靭で高性能な鋼材が求められている。これに加えて当該鋼材を製造するに当たり、省資源且つ省エネルギーであることも重要な課題である。そして、当該鋼材を製造するに当たっては設備を増設ないし新設することなく、しかも従来の製造工程よりも省工程で目的とする鋼材を製造できることが一層望ましい。 In recent years, with the increase in the size of structures and the reduction in weight of automobile parts, stronger and higher performance steel materials are required than ever. In addition to this, it is an important issue to save resources and energy when manufacturing the steel material. And when manufacturing the said steel materials, it is still more desirable to be able to manufacture the target steel materials by a process saving than a conventional manufacturing process, without expanding or newly installing an installation.
従来、主に自動車の車体向けとして高強度で高延性を有し、衝撃エネルギー吸収能にも優れた薄鋼板が多数開発されている。例えば、特許文献1には、高強度と高延性を両立させ、プレス成形性と衝撃エネルギー吸収能に優れた自動車用の冷延鋼板に関する技術が開示されている。これは高価な合金元素の添加量を抑制してフェライト結晶粒の微細化により強度を上昇させ、しかもプレス成形性に重要となる強度と延性とのバランスに優れた薄鋼板である。そしてその製造工程では熱間圧延の後、冷間圧延を行ない、適切な焼鈍を行なうというものである。しかしながら、この技術によれば、MoやNi等の高価な合金元素が少量ではあるが添加必須元素であり、薄鋼板に圧延後、焼鈍処理を必要としている。 Conventionally, many thin steel sheets having high strength and high ductility and excellent impact energy absorption ability have been developed mainly for automobile bodies. For example, Patent Document 1 discloses a technology related to a cold-rolled steel sheet for automobiles that has both high strength and high ductility and is excellent in press formability and impact energy absorption capability. This is a thin steel plate that suppresses the amount of expensive alloy elements added and increases the strength by refining ferrite crystal grains and has an excellent balance between strength and ductility, which is important for press formability. In the manufacturing process, after hot rolling, cold rolling is performed and appropriate annealing is performed. However, according to this technique, a small amount of expensive alloy elements such as Mo and Ni are essential elements to be added, and an annealing treatment is required after rolling into a thin steel plate.
また、非特許文献1には、高価な合金元素を添加せずにMnとSi含有量を高めた0.1%C−5%Mn−2%Siという低炭素鋼に準じる化学成分組成鋼を用い、当該焼鈍後の低温再加熱処理において高含有量のMnにより残留オーステナイトの分率を高めると同時に、高含有量のSiによりフェライト中からオーステナイトへ排出されたCにより残留オーステナイトを安定化させることによる加工硬化指数を高めた鋼板(New TRIP鋼と称される)が開示されている。しかし、このプロセスは薄鋼板に圧延後に複雑なプロセスである焼鈍処理及び低温再加熱処理を必要としており、省エネルギーの観点からの問題が解決されていない。そして、薄鋼板を製造対象鋼材としているので、熱間圧延に加えて冷間圧延工程も必須としている。 Further, Non-Patent Document 1 discloses a chemical composition steel according to a low carbon steel of 0.1% C-5% Mn-2% Si in which the contents of Mn and Si are increased without adding an expensive alloy element. Used to increase the fraction of retained austenite with a high content of Mn in the low-temperature reheating treatment after the annealing, and at the same time stabilize the retained austenite with C discharged from ferrite into austenite with a high content of Si A steel plate (called New TRIP steel) with an increased work hardening index is disclosed. However, this process requires an annealing process and a low-temperature reheating process, which are complicated processes after rolling into a thin steel sheet, and the problem from the viewpoint of energy saving has not been solved. And since the thin steel plate is made into steel material for manufacture, in addition to hot rolling, the cold rolling process is also essential.
一方、製造対象鋼材として薄鋼板を除く構造物等に使用される高強度、高強靭鋼材についても多数開発されている。例えば、特許文献2には、高強度、高延性で、耐遅れ破壊特性に優れ、しかも靭性が飛躍的に向上した高強度鋼材に関する技術が開示されている。この技術によれば、引張強さが1660〜1800MPa、伸び(全伸び)が18.5〜19.2%であって、室温におけるVノッチシャルピー試験の衝撃吸収エネルギーで305〜382J/cm2を有する鋼材が例示されている(特許文献2の表6の実施例1及び実施例17参照)。しかし、この技術においても、化学成分組成として高価格のMoを1.0%程度含有させ、製造工程として、所定の温度及び時間の条件下において焼鈍、焼戻し及び時効処理のいずれかを施した後、350℃以上(AC1−20℃)以下の温度で加工をする(温間加工をする)。 On the other hand, many high-strength and high-tough steel materials used for structures and the like excluding thin steel plates have been developed as steels to be manufactured. For example, Patent Document 2 discloses a technique related to a high strength steel material having high strength, high ductility, excellent delayed fracture resistance, and drastically improved toughness. According to this technique, the tensile strength is 1660 to 1800 MPa, the elongation (total elongation) is 18.5 to 19.2%, and the impact absorption energy of the V-notch Charpy test at room temperature is 305 to 382 J / cm 2 . The steel material which has is illustrated (refer Example 1 and Example 17 of Table 6 of patent document 2). However, even in this technique, about 1.0% of high-priced Mo is contained as a chemical component composition, and after manufacturing, annealing, tempering, or aging treatment is performed under conditions of a predetermined temperature and time. Processing is performed at a temperature of 350 ° C. or higher and (A C1 −20 ° C.) or lower (warm processing is performed).
以上のように、これまでに開示されている技術では省資源、省エネルギーの問題が解決されておらず、また、通常の製造ラインにおいては温間加工を実施するために加工装置に大きな負担を強いることになり、工業的に幅広く利用するには問題がある。 As described above, the technologies disclosed so far have not solved the problems of resource saving and energy saving, and in a normal production line, a heavy burden is imposed on the processing apparatus to perform warm processing. In other words, there is a problem in using industrially widely.
本発明は、以上の点に鑑みて、従来技術では解決することができない上記各種の問題点、即ち、建造物や橋梁等の構造物、自動車の足回り鋼材、機械用歯車等部品に使用される鋼材として、高強度且つ高延性で、衝撃エネルギー吸収能に優れた厚鋼板、形鋼、異形棒鋼、棒鋼及び鋼線等の鋼材を製造するに当たって、高価な合金元素を添加しないで、低炭素鋼の化学成分組成を有する鋼を使用して、製造設備に過大な負荷をかけることなく現有の製造ラインにおいて、多資源・高エネルギーで且つ多工程のために安価かつ所望の鋼材を製造することができない、という問題を、マルテンサイト組織を主体とした冷却速度によらず安定な組織が得られる鋼材を提供することで解決しようとするものである。 In view of the above points, the present invention is used for the above-mentioned various problems that cannot be solved by the prior art, that is, for structures such as buildings and bridges, undercarriage steel materials for automobiles, and mechanical gears. When manufacturing steel materials such as thick steel plates, shaped steels, deformed steel bars, steel bars and steel wires with high strength and high ductility and excellent impact energy absorption capability, low carbon is added without adding expensive alloying elements. Using steel with the chemical composition of steel to produce the desired steel material at low cost for multiple resources, high energy, and multiple processes on the existing production line without overloading the production equipment It is an object of the present invention to provide a steel material that can obtain a stable structure regardless of the cooling rate mainly composed of a martensite structure.
そして製造対象とする鋼材の特性値に関しては、引張強度が1300MPa以上とし、一般的には引張強度の上昇につれて延性が低下するのに対して、本発明では全伸びが18%以上確保され、しかもVノッチシャルピー衝撃試験における上部棚エネルギーが100J/cm2以上確保される高強度非調質鋼材を提供することにある。 And regarding the characteristic value of the steel material to be manufactured, the tensile strength is set to 1300 MPa or more, and generally the ductility decreases as the tensile strength increases, whereas in the present invention, the total elongation is secured to 18% or more. An object of the present invention is to provide a high-strength non-heat treated steel material in which the upper shelf energy in the V-notch Charpy impact test is secured at 100 J / cm 2 or more.
本発明者等は上記の課題を解決するために、鋼材の結晶組織形態の新規組合せの相及びその構成比率と材料特性値との関係を鋭意研究した結果、本発明を完成するに至った。即ち、本発明は以下の特徴を有する。
第1に、化学成分組成が質量%で、C:0.05〜0.20%、Si:1.0〜2.2%、Mn:3.0〜5.5%、Al:0.0005〜0.08%で、残部がFe及び不可避不純物である非調質鋼材であって、組織は、主相がラスマルテンサイトであって、一部に体積分率5%以下の残留オーステナイトからなることを特徴とする強度・延性・靭性に優れた高強度非調質鋼材である。
In order to solve the above problems, the present inventors have intensively studied the phase of a new combination of crystal structure forms of steel materials and the relationship between the constituent ratio and material property values, and as a result, the present invention has been completed. That is, the present invention has the following features.
First, the chemical composition is mass%, C: 0.05 to 0.20%, Si: 1.0 to 2.2%, Mn: 3.0 to 5.5%, Al: 0.0005. in 0.08%, the balance being a non-heat treated steel is Fe and inevitable impurities, organization, main phase a lath martensite, consisting of a volume fraction of 5% or less of retained austenite in a part It is a high strength non-tempered steel material with excellent strength, ductility and toughness.
第2に、上記第1の発明の強度・延性・靭性に優れた高強度非調質鋼材において、主相のラスマルテンサイトが、短径(幅)の平均値が1.0μm以下、長径(長さ)の平均値が7.0μm以下の微細相からなる。 Secondly, in the high-strength non-heat treated steel material excellent in strength, ductility and toughness of the first invention, the main phase of lath martensite has an average value of a short diameter (width) of 1.0 μm or less, a long diameter ( It consists of a fine phase with an average value of (length) of 7.0 μm or less.
第3に、上記第1又は第2の発明の強度・延性・靭性に優れた高強度非調質鋼材において、引張強さが1300MPa以上、全伸びが18.0%以上、2mmVノッチシャルピー衝撃試験による衝撃吸収エネルギーの上部棚エネルギーが100J/cm2以上である。 Thirdly, in the high-strength non-heat treated steel material excellent in strength, ductility and toughness of the first or second invention, the tensile strength is 1300 MPa or more, the total elongation is 18.0% or more, and the 2 mm V notch Charpy impact test. The upper shelf energy of the shock absorption energy due to is 100 J / cm 2 or more.
第4に、上記第1から第3の発明の強度・延性・靭性に優れた高強度非調質鋼材は、1000〜1300℃の範囲内で加熱保持した後、これに塑性相当ひずみeが0.34以上の熱間加工を施すことにより製造される。
Fourth, the high-strength non-heat treated steel material having excellent strength, ductility and toughness according to the first to third inventions is heated and held in the range of 1000 to 1300 ° C., and then the plastic equivalent strain e is 0. Manufactured by performing hot working of 34 or more.
本発明によれば、例えば、1300℃程度に加熱した材料に従来の汎用的な製造設備で、汎用的な塑性加工スケジュール(例えば、熱間圧延スケジュール)で所望の形状・寸法に仕上げ、その間焼入れ・焼戻しや焼鈍といった熱処理工程をしない非調質のままで、上記第3の発明に記載の鋼材特性値である引張強さが1300MPa以上、全伸びが18.0%以上でしかも衝撃吸収エネルギーが2mmVノッチシャルピー衝撃試験片で100J/cm2以上を備えた鋼材を得ることができる。その際、鋼の成分組成としては高価な合金元素は一切不要である。従って、省資源、省エネルギー、低コストで上記鋼材を得ることができる。 According to the present invention, for example, a material heated to about 1300 ° C. is finished with a conventional general-purpose manufacturing facility with a general-purpose plastic working schedule (for example, a hot rolling schedule) to a desired shape and size, and quenched during that time. -The steel material characteristic value described in the third invention is tensile strength of 1300 MPa or more, the total elongation is 18.0% or more, and the impact absorption energy is not tempered without tempering or annealing. A steel material having 100 J / cm 2 or more can be obtained with a 2 mm V notch Charpy impact test piece. At that time, no expensive alloy element is required as the component composition of the steel. Therefore, the steel material can be obtained with resource saving, energy saving and low cost.
以下、本発明に係る鋼材の化学成分、顕微鏡組織及び機械的性質の特徴、並びに当該鋼材の製造方法の特徴について詳細に説明する。
<化学成分組成>
本発明の高強度非調質鋼材における各化学成分組成の範囲は以下の通りである(以下、成分の%はすべて質量%を示す)。
Hereinafter, the characteristics of the chemical composition, microstructure and mechanical properties of the steel material according to the present invention and the characteristics of the method for producing the steel material will be described in detail.
<Chemical component composition>
The range of each chemical component composition in the high-strength non-tempered steel material of the present invention is as follows (hereinafter, “% of components” represents “% by mass”).
C:0.05〜0.20%とする。Cは引張強度を確保するために必要であるが、0.05%未満では本発明に係る鋼材の引張強度を十分に満たさないおそれがあるため、0.05%以上に規定する。一方、0.20%を超えると、鋼材の延性の低下傾向及び溶接性の低下傾向を示すので、上限を0.20%に規定する。 C: Set to 0.05 to 0.20%. C is necessary to ensure the tensile strength, but if it is less than 0.05%, the tensile strength of the steel material according to the present invention may not be sufficiently satisfied, so it is specified to be 0.05% or more. On the other hand, if it exceeds 0.20%, the steel material exhibits a tendency to lower ductility and a tendency to lower weldability, so the upper limit is defined as 0.20%.
Si:1.0〜2.2%とする。Siは本発明における製造工程中の熱間塑性加工後の冷却中に、フェライト中の固溶Cを排出してオーステナイト中に濃化させてオーステナイトを安定化させる。Siのこの作用を十分に発揮させるために1.0%以上が望ましい。一方、Si含有量が過度に高くなると鋼材の加工性が低下傾向を示すので、上限を2.2%とする。 Si: 1.0 to 2.2%. During cooling after hot plastic working during the manufacturing process in the present invention, Si dissolves C in ferrite and concentrates in austenite to stabilize austenite. In order to sufficiently exhibit this effect of Si, 1.0% or more is desirable. On the other hand, if the Si content is excessively high, the workability of the steel material tends to decrease, so the upper limit is made 2.2%.
Mn:3.0〜5.5%とする。Mnは冷却速度によらずに試料の100%マルテンサイト化させるために3%以上とする必要があり、特に第4の発明における製造工程中の熱間塑性加工後の冷却中にマルテンサイト生成作用が発揮される。5.5%を超えると凝固時にMnの偏析が助長するため、製造プロセスが困難となる可能性がある。 Mn: 3.0 to 5.5%. Mn needs to be 3% or more in order to make the sample 100% martensite regardless of the cooling rate, and in particular, the martensite generation action during cooling after hot plastic working during the manufacturing process in the fourth invention. Is demonstrated. If it exceeds 5.5%, segregation of Mn is promoted during solidification, which may make the manufacturing process difficult.
Al:0.0005〜0.08%とする。Alは脱酸のために添加するが、0.0005%未満ではその効果が不十分となる可能性がある。一方、0.08%を超えるとAlNの生成により脆化の問題が起こる可能性がある。 Al: 0.0005 to 0.08%. Al is added for deoxidation, but if it is less than 0.0005%, the effect may be insufficient. On the other hand, if it exceeds 0.08%, the problem of embrittlement may occur due to the formation of AlN.
C、Si、Mn及びAlを除く残部のFe及び不可避不純物に関しては、以下の通りである。工業的に通常行なわれている、例えば、転炉又は電気炉での通常の製鋼精錬工程で残留する溶鋼中の不可避的成分であるP及びSの含有量は、例えば、JIS G3106 溶接構造用圧延鋼材で規定している範囲内、即ち、Pは0.035%以下、Sは0.035%以下であれば望ましい。Pは過度に高含有量であると遅れ破壊特性を劣化させる傾向があり、Sは過度に高含有量であるとMnS介在物が応力集中サイトとなり延性を劣化させる傾向があるので、それぞれ上記含有量以下であれば望ましい。製鋼過程でスクラップ及び雰囲気ガスから混入するその他の混入不可避元素についても、上記に準じた程度の範囲内であればよい。
<顕微鏡組織>
本発明においては、鋼材の顕微鏡組織が、主相がラスマルテンサイトであって、一部、体積分率5%以下の残留オーステナイトを含むことが最大の特徴である。即ち、第2相にはフェライト、ポリゴナルフェライト、準ポリゴナルフェライト、ベイナイト、ベイニティックフェライト、焼戻しマルテンサイト、パーライト及びセメンタイトの内、いずれをも実質的に含んでいない組織である。ここで、実質的に含んでいないとは、倍率10,000倍のSEM及びTEMによる観察でもその存在が確認されないことを意味する。また、上記結晶組織の他には、Nb、Ti、V及びMo等合金元素添加により生成する微細な炭化物系あるいは窒化物系の硬質析出物、並びに球状化セメンタイトも一切認められない組織であることが特徴である。そして、主相がラスマルテンサイトであって、一部、体積分率5%以下の残留オーステナイトを含むことに特徴がある。
The remaining Fe and unavoidable impurities excluding C, Si, Mn and Al are as follows. The contents of P and S, which are inevitable components in the molten steel remaining in a normal steelmaking refining process in, for example, a converter or an electric furnace, which are usually performed industrially, are, for example, JIS G3106 Rolling for welded structures It is desirable that the steel is within the range defined by the steel material, that is, P is 0.035% or less and S is 0.035% or less. If P is excessively high, the delayed fracture characteristics tend to deteriorate. If S is excessively high, MnS inclusions tend to become stress concentration sites and deteriorate ductility. Less than the amount is desirable. Other mixed inevitable elements mixed from scrap and atmospheric gas in the steelmaking process may be within a range similar to the above.
<Microscopic structure>
In the present invention, the greatest feature of the microstructure of the steel material is that the main phase is lath martensite and partly contains retained austenite having a volume fraction of 5% or less. That is, the second phase is a structure that substantially does not contain any of ferrite, polygonal ferrite, quasi-polygonal ferrite, bainite, bainitic ferrite, tempered martensite, pearlite, and cementite. Here, “substantially free” means that the presence is not confirmed even by observation with a 10,000 times SEM and TEM. In addition to the above crystal structure, fine carbide-based or nitride-based hard precipitates formed by addition of alloy elements such as Nb, Ti, V, and Mo, and spheroidized cementite are not observed at all. Is a feature. The main phase is lath martensite, and partly contains residual austenite having a volume fraction of 5% or less.
これまでに、高強度、高延性で衝撃吸収エネルギー能に優れた鋼材の開示は多数あるが、上述した顕微鏡組織を有する鋼材は知見されていない。そして上記組織であって、更に、上記ラスマルテンサイトの短径(幅)の平均値が1.0μm以下で、長径(長さ)の平均値が7.0μm以下という微細であれば、強度−延性バランスに優れた鋼材を得る観点から一層望ましい。
<製造方法>
次に、本発明の鋼材を得るための好ましい製造方法を説明する。
(熱間加工に至るまでの製造方法)
素材は通常方法で形さえ整えられれば良い。
(熱間加工条件)
上記で得られた素材の熱間加工方式としては、工業的に行なわれている厚鋼板製造ラインにおける平ロール圧延、極厚鋼板製造ラインにおける鍛造、棒鋼又は鋼線材製造ラインにおける溝ロール圧延、及び条鋼又は形鋼製造ラインにおける形ロール圧延の内のいずれであってもよい。これらいずれかの加工方式により、素材に対して所要の塑性相当ひずみを与える。
To date, there are many disclosures of steel materials having high strength, high ductility, and excellent impact absorption energy capability, but no steel materials having the above-described microstructure have been found. And if it is the said structure, and the average value of the short diameter (width) of the lath martensite is 1.0 μm or less and the average value of the long diameter (length) is 7.0 μm or less, the strength − It is more desirable from the viewpoint of obtaining a steel material having an excellent ductility balance.
<Manufacturing method>
Next, the preferable manufacturing method for obtaining the steel material of this invention is demonstrated.
(Manufacturing method up to hot working)
The material only needs to be shaped in the usual way.
(Hot processing conditions)
As the hot working method of the raw material obtained above, flat roll rolling in a thick steel plate production line, industrial forging, forging in a very thick steel plate production line, groove roll rolling in a bar or steel wire production line, and Any of the shape roll rolling in a steel bar or a shape steel production line may be sufficient. By any one of these processing methods, a required plastic equivalent strain is given to the material.
上記の加工方式により、素材に導入される圧縮ひずみとせん断ひずみの入り方は異なる。そこで、全応力成分や全ひずみ成分の量や分布に関して理論的に塑性ひずみを算出する方法として、有限要素法(finite element methode:FEM)がある。塑性ひずみの計算については、参考文献(春海佳三郎、他「有限要素法入門」(共立出版(株):1990年3月15日)に詳述されている。しかしここでは、工業的に簡便に用いることができる塑性相当ひずみを用いてもよい。有限要素法計算で得られる塑性ひずみを用いれば一層望ましいが、ここでは工業的に簡便な、下記式(1)で定義される塑性相当ひずみ(e)を塑性ひずみの指標とする。 Depending on the above processing method, the way of entering the compressive strain and shear strain introduced into the material is different. Therefore, there is a finite element method (FEM) as a method for theoretically calculating the plastic strain with respect to the amount and distribution of the total stress component and the total strain component. The calculation of plastic strain is described in detail in the reference (Kasaburo Harumi, et al. “Introduction to Finite Element Method” (Kyoritsu Shuppan Co., Ltd .: March 15, 1990). The plastic equivalent strain that can be used for the plastic strain may be used, but it is more desirable to use the plastic strain obtained by the finite element method calculation, but here, the plastic equivalent strain defined by the following formula (1) is industrially simple. Let (e) be an index of plastic strain.
e=−ln(1−R/100)・・・・・・(1)
但し、Rは減面率(%)であり、素材のC方向断面積をS0とし、熱間加工後のC方向断面積をSとすると、下記式(2)で表される。
e = -ln (1-R / 100) (1)
However, R is an area reduction rate (%), and when the C direction sectional area of the material is S 0 and the C direction sectional area after hot working is S, it is expressed by the following formula (2).
R={(S0−S)/S0 }×100・・・・・・(2)
そして、熱間加工工程で素材に導入すべき所要の塑性相当ひずみ(e)は、後述する実施例1の実験での1200℃で5分間の再加熱後の断面45mm角から38mm角への鍛造加工による塑性相当ひずみ(e)の量に基づき、e≧0.34と規定する。これにより、前記<顕微鏡組織>の項で述べた通りの、主相がラスマルテンサイトであって、一部、体積分率5%以下の残留オーステナイトを含む顕微鏡組織を有する鋼材を得ることができる。
R = {(S 0 −S) / S 0 } × 100 (2)
The required plastic equivalent strain (e) to be introduced into the raw material in the hot working step is forging from a 45 mm square to a 38 mm square cross section after reheating at 1200 ° C. for 5 minutes in the experiment of Example 1 described later. Based on the amount of plastic equivalent strain (e) due to processing, it is defined as e ≧ 0.34. Thereby, as described in the section of <Microstructure>, it is possible to obtain a steel material having a microstructure including a retained austenite whose main phase is lath martensite and a volume fraction of 5% or less. .
次に、加工温度及びパススケジュールについて述べる。加工温度はオーステナイト温度領域の中でも、現実の操業性に有利な約1000〜1300℃、好ましくは1000〜1200℃の温度範囲内において約20〜60分間加熱保持してから熱間加工を開始する。熱間加工方式としては、工業的に行なわれている厚鋼板製造ラインにおける平ロール圧延、極厚鋼板製造ラインにおける鍛造、棒鋼又は鋼線材製造ラインにおける溝ロール圧延、及び条鋼又は形鋼製造ラインにおける形ロール圧延の内のいずれであってもよい。 Next, the processing temperature and the pass schedule will be described. Within the austenite temperature range, the processing temperature is about 1000 to 1300 ° C., preferably 1000 to 1200 ° C., which is advantageous for actual operability, and is heated and held for about 20 to 60 minutes, and then hot processing is started. Hot working methods include flat roll rolling in industrial steel plate production lines, forging in extra-thick steel plate production lines, groove roll rolling in steel bar or steel wire production lines, and strip or shape steel production lines. Any of the shape roll rolling may be used.
仕上げ温度は、後述する図3の実施例2における鍛造終了温度(688℃)以上であれば良い。この温度以上とした理由は、本発明鋼材の組織形態の特徴(主相がラスマルテンサイトであって、一部、体積分率5%以下の残留オーステナイトを含む)を確保するためである。但し、マルテンサイト変態開始温度が約350℃のため、学術的には350℃以上であればよいが、実際の工業生産においては、本発明の課題の1つでもある生産設備に過剰な負荷をかけないという点から、688℃以上であれば良い。 The finishing temperature should just be more than the forge completion temperature (688 degreeC) in Example 2 of FIG. 3 mentioned later. The reason why the temperature is higher than this temperature is to ensure the characteristics of the structure of the steel of the present invention (the main phase is lath martensite and partially contains residual austenite having a volume fraction of 5% or less). However, since the martensitic transformation start temperature is about 350 ° C., it may be 350 ° C. or higher academically. However, in actual industrial production, an excessive load is applied to the production equipment which is one of the problems of the present invention. In view of not applying, it may be 688 ° C. or higher.
更に、仕上げ後の冷却速度は、以下の予備実験を行って、空冷、ガス冷却あるいは水冷のいずれであってもよいことを見出した。
<予備実験>
本発明者等は、0.10%C−2.0%Si−5.0%Mnなる低炭素鋼で断面38mm角×400mm長さの試験材を用い、1200℃で1時間加熱保持した後、圧延開始温度と圧延後の冷却方法を変えた以下の6種の方法で溝ロール圧延をして、いずれも断面14.3mm角に仕上げた。表1に溝ロール圧延のパススケジュールを示す。
(1)1200℃で1時間加熱保持した後、溝ロール圧延を開始して断面14.3mm角に仕上げた。仕上げ温度は715℃であった。その後、水冷した場合と空冷した場合の2種の試験材を調製した。
(2)1200℃で1時間加熱保持した後、900℃まで空冷した後、溝ロール圧延を開始して断面14.3mm角に仕上げた。仕上げ温度は685℃であった。その後、水冷した場合と空冷した場合の2種の試験材を調製した。
(3)1200℃で1時間加熱保持した後、600℃まで空冷した後、溝ロール圧延を開始して断面14.3mm角に仕上げた。仕上げ温度は610℃であった。その後、水冷した場合と空冷した場合の2種の試験材を調製した。
Furthermore, the cooling rate after finishing was found to be air cooling, gas cooling or water cooling by conducting the following preliminary experiment.
<Preliminary experiment>
The present inventors used a test material having a cross section of 38 mm square x 400 mm length in a low carbon steel of 0.10% C-2.0% Si-5.0% Mn, and heated and held at 1200 ° C for 1 hour. Then, groove rolling was performed by the following six methods with different rolling start temperatures and cooling methods after rolling, and all were finished to have a cross section of 14.3 mm square. Table 1 shows the pass schedule for groove roll rolling.
(1) After heating and holding at 1200 ° C. for 1 hour, groove roll rolling was started to finish a cross section of 14.3 mm square. The finishing temperature was 715 ° C. Thereafter, two types of test materials were prepared for water cooling and air cooling.
(2) After heating and holding at 1200 ° C. for 1 hour, after air cooling to 900 ° C., groove roll rolling was started to finish a cross section of 14.3 mm square. The finishing temperature was 685 ° C. Thereafter, two types of test materials were prepared for water cooling and air cooling.
(3) After heating and holding at 1200 ° C. for 1 hour, after air cooling to 600 ° C., groove roll rolling was started to finish a cross section of 14.3 mm square. The finishing temperature was 610 ° C. Thereafter, two types of test materials were prepared for water cooling and air cooling.
以上6種の試験材について、ビッカース硬さ(試験荷重:10kg)を各5点ずつ測定した。その結果を表2に示す。 With respect to the above six types of test materials, the Vickers hardness (test load: 10 kg) was measured at five points each. The results are shown in Table 2.
表2に示すように、6種の試料のすべてにおいて、ビッカース硬さはほぼ同一レベルのHV(10kg荷重試験)=465〜500の範囲内にあることが確認された。 As shown in Table 2, it was confirmed that the Vickers hardness was in the range of almost the same level of H V (10 kg load test) = 465 to 500 in all of the six types of samples.
以上により、硬さが強度と比例する関係から、熱間圧延における圧延温度、及び圧延終了後の冷却速度によらず、ほぼ一定の強度特性の材料が得られることがわかった。これにより、素材の熱間加工温度に変動があっても一定の材質特性が保証できることがわかる。加工温度領域が大きく異なっていても、鋼材の特性に差が生じない理由は、前述した通り、マルテンサイト変態開始温度が約350℃と低いため、これよりも高温領域であれば加工温度の影響は生じなかったとものと考えられる。
<本発明の鋼材の特性>
上記のようにして得られた熱間加工鋼材、即ち本発明に係る鋼材は、その用途により適宜、表面のスケール除去等の表面処理は施すものの、鋼材の機械的性質を向上させるための調質処理は一切しないものとする。非調質のままで、本発明が目標としている機械的性質を全て満たしているからである。即ち、引張強さが1300MPa以上、全伸びが18.0%以上でしかも衝撃吸収エネルギーが2mmVノッチシャルピー衝撃試験片で上部棚エネルギーが100J/cm2以上を有する。これは、本発明の鋼材が、主相がラスマルテンサイトであって、一部、体積分率5%以下の残留オーステナイトを含みその短径及び長径が比較的微細な結晶粒径を有するからである。
From the above, it was found that a material having substantially constant strength characteristics can be obtained regardless of the rolling temperature in hot rolling and the cooling rate after the end of rolling because the hardness is proportional to the strength. As a result, it can be seen that a certain material characteristic can be guaranteed even if the hot working temperature of the material varies. The reason why there is no difference in the characteristics of the steel material even if the processing temperature region is greatly different is that the martensite transformation start temperature is as low as about 350 ° C as described above. It is thought that the problem did not occur.
<Characteristics of the steel material of the present invention>
The hot-worked steel material obtained as described above, that is, the steel material according to the present invention is subjected to surface treatment such as surface scale removal depending on its use, but is tempered to improve the mechanical properties of the steel material. No processing shall be performed. This is because all the mechanical properties targeted by the present invention are satisfied while remaining untempered. That is, the tensile strength is 1300 MPa or more, the total elongation is 18.0% or more, and the impact absorption energy is 2 mmV notch Charpy impact test piece and the upper shelf energy is 100 J / cm 2 or more. This is because the steel material of the present invention has a lath martensite main phase, and partially contains residual austenite with a volume fraction of 5% or less, and the minor axis and major axis have a relatively fine crystal grain size. is there.
以下、実施例により本発明を更に具体的に説明する。なお、本発明は、下記実施例によって制限されず、前記及び後記の趣旨に適合し得る範囲で適切な改変を行って実施することも可能であり、これらはいずれも本発明の技術的範囲内に含まれる。
<実施例1>
本発明の実施例1の鋼材を下記の熱間鍛造加工により得た。
Hereinafter, the present invention will be described more specifically with reference to examples. It should be noted that the present invention is not limited by the following examples, and can be implemented with appropriate modifications within a range that can be adapted to the gist described above and below, both of which are within the technical scope of the present invention. include.
<Example 1>
The steel material of Example 1 of the present invention was obtained by the following hot forging process.
先ず、溶解用主原料として電解鉄、電解Mn及び金属Siを使用し、高周波真空誘導溶解炉を用いて、表3に示す化学成分組成の溶鋼(単位:質量%)を溶製し、縦95mm×横95mm×高さ450mmの鋼塊を鋳造した。これを鋼塊番号1とする。この鋼塊の上部側半分を切り出して素材とした。 First, electrolytic iron, electrolytic Mn, and metal Si are used as main raw materials for melting, and molten steel (unit: mass%) having chemical composition shown in Table 3 is melted using a high-frequency vacuum induction melting furnace. A steel ingot having a width of 95 mm and a height of 450 mm was cast. This is steel ingot number 1. The upper half of this steel ingot was cut out and used as a material.
上記素材を加熱昇温し、1200℃で1時間加熱保持した後、縦95mm×横95mmの角形状断面の素材に、縦と横とを交互に1回ずつのプレス鍛造を表4に示す圧下のスケジュールで行い、縦45mm×横45mmの角形状断面(以後、45mm角という)まで鍛造した。この時点の温度は約800℃であり、ここから再加熱昇温して1200℃で5分間加熱保持した後、更に鍛造を施して38mm角とし、そして最後に材料全体の形状を直線状に矯正するために極軽鍛造を施して、断面が38mm角の棒状鋼材に仕上げた。鍛造仕上温度は1000℃と推定され、鍛造終了後直ちに空冷し、室温まで冷却した。 After heating and heating the above materials for 1 hour at 1200 ° C., press forging is performed as shown in Table 4 in which the vertical and horizontal press forges are alternately applied to the material having a square cross section of 95 mm long × 95 mm wide. The forging was performed to a square cross section (hereinafter referred to as 45 mm square) of 45 mm length × 45 mm width. The temperature at this point is about 800 ° C., reheated from here, heated and held at 1200 ° C. for 5 minutes, further forged to 38 mm square, and finally straightened the overall shape of the material In order to achieve this, extremely light forging was applied to finish a rod-shaped steel material having a cross section of 38 mm square. The forging finishing temperature was estimated at 1000 ° C., and immediately after the forging was completed, the product was air cooled and cooled to room temperature.
上記鍛造スケジュール並びに塑性相当ひずみ(e)及び減面率(R)を、表4に示す。 Table 4 shows the forging schedule, plastic equivalent strain (e), and area reduction rate (R).
ここで、塑性相当ひずみ(e)及び減面率(R)は、前記<製造方法>(熱間加工条件)の項で述べたものと同じであり、それぞれ下記式(1)及び(2)で算出した。 Here, the plastic equivalent strain (e) and the area reduction ratio (R) are the same as those described in the section of <Manufacturing method> (Hot working conditions), and the following formulas (1) and (2) respectively. Calculated with
e=−ln(1−R/100) ・・・・・・・(1)
R=(S0−S)/S0 ×100 ・・・・・・・(2)
上記のようにして本実施例1の鋼材を得た。
e = -ln (1-R / 100) (1)
R = (S 0 −S) / S 0 × 100 (2)
The steel material of Example 1 was obtained as described above.
なお、後述する比較例1の試験では、本実施例1で得られた鋼材から確性試験用として必要な試験片を採取した後の残鋼材を用いて、これに更に、後述するように焼鈍処理を施すことによって、比較例1の鋼材を得た。
<実施例2>
本発明の実施例2の鋼材を下記の熱間鍛造加工により得た。
In the test of Comparative Example 1 to be described later, using the remaining steel material after collecting a test piece necessary for the reliability test from the steel material obtained in Example 1, an annealing treatment is further performed as described later. The steel material of the comparative example 1 was obtained by giving.
<Example 2>
The steel material of Example 2 of the present invention was obtained by the following hot forging process.
実施例1に準じて溶解用主原料として電解鉄、電解Mn及び金属Siを使用し、高周波真空誘導溶解炉を用いて、表5に示す化学成分組成の溶鋼(単位:質量%)を溶製し、縦95mm×横95mm×高さ450mmの鋼塊を鋳造した。これを鋼塊番号2とする。この鋼塊の上部側半分を切り出して素材とした。 In accordance with Example 1, electrolytic iron, electrolytic Mn and metal Si were used as the main raw materials for melting, and molten steel (unit: mass%) having the chemical composition shown in Table 5 was melted using a high-frequency vacuum induction melting furnace. A steel ingot having a length of 95 mm, a width of 95 mm, and a height of 450 mm was cast. This is steel ingot number 2. The upper half of this steel ingot was cut out and used as a material.
上記素材を加熱昇温し、1200℃で1時間加熱保持した後、縦95mm×横95mmの角形状断面の素材に、縦と横とを交互に1回ずつの鍛造を施して、38mm角まで鍛造した。この間途中で再加熱することなく鍛造をした。そして最後に材料全体の形状を直線状に矯正するために極軽鍛造を行なって、断面が38mm角の棒状鋼材に仕上げた。鍛造仕上温度は688℃であり、鍛造が終了後直ちに空冷し、室温まで冷却した。 The above material is heated and heated at 1200 ° C. for 1 hour, and then forged once alternately in the vertical and horizontal directions to a square cross section of 95 mm long by 95 mm wide to 38 mm square Forged. During this time, forging was performed without reheating in the middle. Finally, extremely light forging was performed to correct the shape of the entire material into a straight line, and a rod-shaped steel material having a cross section of 38 mm square was finished. The forging finishing temperature was 688 ° C., and after forging was completed, the product was air-cooled immediately and cooled to room temperature.
上記鍛造スケジュール並びに塑性相当ひずみ(e)及び減面率(R)を表6に示す。 Table 6 shows the forging schedule, plastic equivalent strain (e), and area reduction rate (R).
冷却中の相変態をフォーマスター試験で熱膨張測定により決定し、そこからCCT図を作成したものを図1に示す。図1から冷却速度が67K/sec〜0.5K/secの範囲で、Ms点が〜340℃であることが分かる。 The phase transformation during cooling is determined by thermal expansion measurement in a four master test, and a CCT diagram is created from the result shown in FIG. It can be seen from FIG. 1 that the cooling rate is 67 K / sec to 0.5 K / sec and the Ms point is ˜340 ° C.
通常の鋼では、冷却速度が遅いと図示しないγ/α変態のノーズに当たり、フェライトが生成する。しかしながら、本発明の鋼では、γ/α変態のフェライト生成範囲が成分元素の効果で高時間側に遷移したため、冷却速度が0.5K/secと遅い場合でもγ/α変態のノーズに当たらずにMs点に当たり、マルテンサイト変態が起こる。 In normal steel, when the cooling rate is low, ferrite hits a nose of γ / α transformation (not shown) and ferrite is generated. However, in the steel of the present invention, the ferrite generation range of the γ / α transformation has shifted to the high time side due to the effect of the component elements, so even when the cooling rate is as low as 0.5 K / sec, it does not hit the nose of the γ / α transformation. At the Ms point, martensitic transformation occurs.
このように、本発明材は組織の冷却速度依存性が低いことがわかる。また、実施例1の材料の組織をEBSDで決定したものを図2に示す。これから組織の95%以上はマルテンサイト(α’)相であり、残り5%以下のオーステナイト(γ)相から成っていることが決定される。 Thus, it can be seen that the material of the present invention has a low dependence on the cooling rate of the tissue. Moreover, what determined the structure | tissue of the material of Example 1 by EBSD is shown in FIG. From this, it is determined that 95% or more of the structure is a martensite (α ′) phase and the remaining 5% or less is an austenite (γ) phase.
なお、図3に示すように鍛造は1200℃で開始し、終了時の温度は688℃であった。その後、空冷を行ない、約200℃に至るまでの材料の温度変化を示す。
<実施例3>
本発明の実施例3の鋼材を下記の熱間鍛造加工により得た。
As shown in FIG. 3, forging started at 1200 ° C., and the temperature at the end was 688 ° C. Then, air cooling is performed and the temperature change of the material until it reaches about 200 degreeC is shown.
<Example 3>
The steel material of Example 3 of the present invention was obtained by the following hot forging process.
実施例2で溶製され、鋳造された鋼塊(鋼塊番号2)の下部側半分を素材とした。即ち、溶解用主原料として電解鉄、電解Mn及び金属Siを使用し、高周波真空誘導溶解炉を用いて、表5に示した化学成分組成の溶鋼(単位:質量%)を溶製し、縦95mm×横95mm×高さ450mmの鋼塊(鋼塊番号2)を鋳造した。この鋼塊の下部側半分を切り出して素材とした。 The lower half of the steel ingot (steel ingot number 2) melted and cast in Example 2 was used as the material. That is, using electrolytic iron, electrolytic Mn, and metal Si as the main raw materials for melting, using a high-frequency vacuum induction melting furnace, molten steel (unit: mass%) having the chemical composition shown in Table 5 was melted, A steel ingot (steel ingot number 2) of 95 mm × width 95 mm × height 450 mm was cast. The lower half of the steel ingot was cut out and used as a material.
上記素材を加熱昇温し、1200℃で1時間加熱保持した後、縦95mm×横95mmの角形状断面の素材に、縦と横とを交互に1回ずつの鍛造を施して、38mm角まで鍛造した。この間途中で再加熱することなく鍛造をした。そして最後に材料全体の形状を直線状に矯正するために極軽鍛造を行なって、断面が38mm角の棒状鋼材に仕上げた。鍛造仕上温度は690℃であり、鍛造が終了後直ちに水冷し、室温まで冷却した。 The above material is heated and heated at 1200 ° C. for 1 hour, and then forged once alternately in the vertical and horizontal directions to a square cross section of 95 mm long by 95 mm wide to 38 mm square Forged. During this time, forging was performed without reheating in the middle. Finally, extremely light forging was performed to correct the shape of the entire material into a straight line, and a rod-shaped steel material having a cross section of 38 mm square was finished. The forging finishing temperature was 690 ° C., and immediately after the forging was completed, the product was cooled with water and cooled to room temperature.
上記鍛造スケジュール並びに塑性相当ひずみ(e)及び減面率(R)は、実施例2における設定条件と同一である。即ち、表6に示した通りである。これよりわかるように、実施例3と実施例2との相違点は、鋼材の仕上目標寸法(38mm角)まで鍛造した後に、実施例2においては空冷したのに対して、実施例3においては水冷したことだけである。 The forging schedule, plastic equivalent strain (e), and area reduction rate (R) are the same as the setting conditions in Example 2. That is, it is as shown in Table 6. As can be seen, the difference between Example 3 and Example 2 is that, after forging to the finishing target size (38 mm square) of the steel material, it was air-cooled in Example 2, whereas in Example 3, It is only water-cooled.
次に、本願発明の範囲外である比較例について説明する。
<比較例1>
比較例1の鋼材は、前述した実施例1で得られた鋼材(鍛造加工終了ままの鋼材)からその確性試験用の試験片を採取した後の残鋼材を、更に、550℃で1時間加熱保持の焼鈍処理を行ない、その後室温まで空冷することにより得た。従って、素材の化学成分組成は前記表3に示した通りであり、この素材を表4に示した鍛造スケジュール及び塑性相当ひずみの条件で鍛造し、その後で上記550℃で1時間加熱保持の焼鈍処理をした。
Next, a comparative example that is outside the scope of the present invention will be described.
<Comparative Example 1>
The steel material of Comparative Example 1 was further heated at 550 ° C. for 1 hour after collecting the test piece for the test of accuracy from the steel material obtained in Example 1 (the steel material after completion of the forging process). It was obtained by carrying out a holding annealing treatment and then cooling to room temperature. Therefore, the chemical composition of the material is as shown in Table 3, and the material is forged under the forging schedule and plastic equivalent strain conditions shown in Table 4, and then annealed at 550 ° C. for 1 hour. Processed.
このように、最後に焼鈍処理を施した試験を行なった理由は、従来の高強度構造用鋼材の製造技術によれば、熱間圧延等の加工をした後、適宜焼入れをし、次いで焼戻し処理を施すことにより延性と靭性の回復を図ることが行なわれるので、本試験においても同じ目的のために焼鈍処理を施した。しかしながら、本比較例1においては、実施例1で得られた鋼材を更に焼鈍処理を施すと、延性は若干向上するものの、靭性が極端に低下した。この焼鈍処理で得られたような靭性の劣化は、これまでの報告には見当たらない。
<比較例2>
比較例2の鋼材を次のようにして得た。
As described above, the reason why the last annealing treatment was performed is that, according to the conventional manufacturing technology for high strength structural steel, after hot rolling or the like, it is appropriately quenched and then tempered. In order to recover ductility and toughness, the annealing treatment was performed for the same purpose in this test. However, in this Comparative Example 1, when the steel material obtained in Example 1 was further subjected to annealing treatment, the ductility was slightly improved, but the toughness was extremely lowered. The deterioration of toughness as obtained by this annealing treatment has not been found in previous reports.
<Comparative example 2>
The steel material of Comparative Example 2 was obtained as follows.
溶解用主原料として電解鉄、電解Mn及び金属Siを使用し、高周波真空誘導溶解炉を用いて、表7に示す化学成分組成の溶鋼(単位:質量%)を溶製し、縦95mm×横95mm×高さ450mmの鋼塊を鋳造した。これを鋼塊番号3とする。この鋼塊の上部側半分を切り出して素材とした。 Using electrolytic iron, electrolytic Mn and metallic Si as the main raw materials for melting, using a high-frequency vacuum induction melting furnace, the molten steel (unit: mass%) having the chemical composition shown in Table 7 is melted, and the length is 95 mm × width A steel ingot measuring 95 mm × 450 mm in height was cast. This is steel ingot number 3. The upper half of this steel ingot was cut out and used as a material.
上記素材を加熱昇温し、1200℃で1時間加熱保持した後、縦95mm×横95mmの角形状断面の素材に、縦と横とを交互に1回ずつの鍛造を施して、38mm角まで鍛造した。この間途中で再加熱することなく鍛造をした。そして最後に材料全体の形状を直線状に矯正するために極軽鍛造を行なって、断面が38mm角の棒状とした。鍛造スケジュール並びに塑性相当ひずみ(e)及び減面率(R)は、表6に示した実施例2及び実施例3と同じ条件であり、鍛造終了温度は680℃であり、鍛造終了後は実施例2と同じで直ちに空冷し、室温まで冷却した。 The above material is heated and heated at 1200 ° C. for 1 hour, and then forged once alternately in the vertical and horizontal directions to a square cross section of 95 mm long by 95 mm wide to 38 mm square Forged. During this time, forging was performed without reheating in the middle. Finally, extremely light forging was performed in order to straighten the shape of the entire material into a bar shape with a cross section of 38 mm square. The forging schedule, plastic equivalent strain (e), and area reduction rate (R) are the same conditions as in Example 2 and Example 3 shown in Table 6, the forging end temperature is 680 ° C., and after forging is completed. As in Example 2, it was immediately air cooled and cooled to room temperature.
更に、この38mm角の棒材を、表8に示すように、下記の温間温度域における溝ロール圧延により、14.3mm角の棒材とした。38mm角から14.3mm角までの溝ロール圧延条件は、表8に示すように、650℃に昇温加熱して1時間加熱保持した後、溝ロールによる3パス圧延後に650℃の保持炉で5分間加熱保持する工程を3回行ない、第パス圧延後に、水冷した。 Further, as shown in Table 8, this 38 mm square bar was formed into a 14.3 mm square bar by groove roll rolling in the following warm temperature range. As shown in Table 8, the groove roll rolling conditions from 38 mm square to 14.3 mm square were heated to 650 ° C. and heated for 1 hour, and then held in a holding furnace at 650 ° C. after 3-pass rolling with a groove roll. The process of heating and holding for 5 minutes was performed 3 times, and after the first pass rolling, water cooling was performed.
その後の熱処理として、675℃で2分間の加熱保持後にHeガス中で室温まで冷却し(焼鈍処理)、次いで400℃まで昇温して5分間の加熱保持後にHeガス中で室温まで冷却する(過時効処理)という2回の熱処理を施した。こうして比較例2の鋼材を得た。 As a subsequent heat treatment, after heating and holding at 675 ° C. for 2 minutes, the mixture is cooled to room temperature in He gas (annealing treatment), then heated to 400 ° C. and held for 5 minutes and then cooled to room temperature in He gas ( Two heat treatments (overaging treatment) were performed. Thus, a steel material of Comparative Example 2 was obtained.
以下、実施例1、2及び3、並びに比較例1及び2の各鋼材について、顕微鏡組織試験、硬さ試験、引張試験及びシャルピー衝撃試験を行なって、材質特性水準を明らかにした。また、フォーマスター試験により、各実施例及び比較例と実質的に同一化学成分組成の鋼材のAc1変態点、並びに実施例1、2及び比較例1において鍛造加工の末期又は終了後の空冷過程における冷却時の変態温度を測定した。
<試験方法>
(顕微鏡組織試験)
顕微鏡組織試験は、各鋼材の圧延方向に垂直な断面(C方向断面)の中心部について、走査型電子顕微鏡(SEM)で金属組織を観察した。
(硬さ試験)
試験荷重10gのマイクロビッカース試験により、2相組織の各結晶組織内の硬さを測定し、また試験荷重10kgのビッカース試験により、2相組織全体の平均硬さを測定した。
(引張試験)
引張試験は、比較例2以外の鋼材からL方向の丸棒引張試験片(試験部分の平行部直径が3.5mmφ、長さが24.5mm)を調製して、引張強度、降伏強度、全伸び及び絞りを測定した。比較例2
についてはL方向の微小引張試験片を用い、フォーマスター装置により応力−ひずみ曲線を測定した。
(シャルピー衝撃試験)
シャルピー衝撃試験は、各鋼材からL方向のJIS Z2242のシャルピー衝撃試験方法により、標準試験片(10×10×55mm、2mmVノッチ)を調製し、−196℃〜100℃の衝撃試験を行い、上部棚エネルギー及び0℃における吸収エネルギーを測定し、破面観察により、脆性破面遷移温度を求めた。
<特性試験結果>
図4〜図8に、実施例1〜3及び比較例1、2の顕微鏡組織試験によるSEM写真を示し、図9〜図13に、実施例1〜3及び比較例1、2の引張試験による応力−ひずみ曲線
を示す。
Hereinafter, for each of the steel materials of Examples 1, 2, and 3 and Comparative Examples 1 and 2, a microstructural test, a hardness test, a tensile test, and a Charpy impact test were performed to clarify the material property level. Further, in the Formaster test, the Ac1 transformation point of the steel material having substantially the same chemical composition as each Example and Comparative Example, and in the air cooling process after the end of the forging process or after completion in Examples 1, 2 and Comparative Example 1 The transformation temperature during cooling was measured.
<Test method>
(Microscopic examination)
In the microscopic structure test, the metal structure was observed with a scanning electron microscope (SEM) at the center of a cross section (C direction cross section) perpendicular to the rolling direction of each steel material.
(Hardness test)
The hardness in each crystal structure of the two-phase structure was measured by a micro Vickers test with a test load of 10 g, and the average hardness of the entire two-phase structure was measured by a Vickers test with a test load of 10 kg.
(Tensile test)
The tensile test was performed by preparing a round bar tensile test piece in the L direction (a parallel part diameter of 3.5 mmφ and a length of 24.5 mm) from a steel material other than Comparative Example 2, and tensile strength, yield strength, total Elongation and drawing were measured. Comparative Example 2
Was measured for stress-strain curves using a four-master apparatus using a micro tensile specimen in the L direction.
(Charpy impact test)
The Charpy impact test is a standard test piece (10 × 10 × 55 mm, 2 mmV notch) prepared from each steel material by the JIS Z2242 Charpy impact test method in the L direction, and an impact test of −196 ° C. to 100 ° C. is performed. The shelf energy and the absorbed energy at 0 ° C. were measured, and the brittle fracture surface transition temperature was determined by observing the fracture surface.
<Characteristic test results>
FIGS. 4 to 8 show SEM photographs of the microstructure tests of Examples 1 to 3 and Comparative Examples 1 and 2, and FIGS. 9 to 13 show the tensile tests of Examples 1 to 3 and Comparative Examples 1 and 2. A stress-strain curve is shown.
表10に、実施例1〜3及び比較例1、2の引張試験及びシャルピー衝撃試験による測定値を示す。そして、表10には、本発明鋼材としての合否判定基準値と、開発目標値とを併記した。
In Table 10 , the measured value by the tensile test and Charpy impact test of Examples 1-3 and Comparative Examples 1 and 2 is shown. And in Table 10 , the acceptance criteria value as this invention steel material and the development target value were written together.
これよりわかるように、実施例1〜実施例3は全て、本発明鋼材としての金属組織とその構成比率が本発明鋼材として合格している。即ち、主相がラスマルテンサイトであって、一部、体積分率5%以下の残留オーステナイトを含んでいる。そして、引張強度(TS)は1353MPa以上であり、全伸び(El)は18.0%以上となっている。また、引張強度(TS)×全伸び(El)は26406MPa・%以上となっており、強度−延性バランスに優れている。更に、衝撃吸収エネルギーに関しては、上部棚エネルギー(USE)が175J/cm2以上と優れている。以上の結果として、材質特性指数を引張強度(TS)×全伸び(El)×上部棚エネルギー(USE)で表わすと、TS×El×USEが4621×103MPa・%・J/cm2以上となり、優れていることが分かる。 As can be seen from this, all of Examples 1 to 3 have passed the metal structure as the steel of the present invention and the component ratio thereof as the steel of the present invention. That is, the main phase is lath martensite and partly contains residual austenite with a volume fraction of 5% or less. The tensile strength (TS) is 1353 MPa or more, and the total elongation (El) is 18.0% or more. Further, the tensile strength (TS) × total elongation (El) is 26406 MPa ·% or more, which is excellent in the strength-ductility balance. Furthermore, regarding the shock absorption energy, the upper shelf energy (USE) is excellent at 175 J / cm 2 or more. As a result of the above, when the material property index is expressed by tensile strength (TS) × total elongation (El) × upper shelf energy (USE), TS × El × USE is 4621 × 10 3 MPa ·% · J / cm 2 or more. It turns out that it is excellent.
なお、主相の硬さ(HV)は459〜510の範囲内にあり、これらの硬さ測定値により、上記金属組織種とその組織相の構成分率の測定値との妥当性を、引張強度(TS)とビッカース硬さ(HV)との関係からも確認することができる。 In addition, the hardness (H V ) of the main phase is in the range of 459 to 510, and the validity of the measured values of the metal structure type and the constituent fraction of the texture phase is determined by these hardness measurement values. This can also be confirmed from the relationship between the tensile strength (TS) and the Vickers hardness (H V ).
次に、熱間鍛造による仕上がり鋼材への塑性相当ひずみ量をみると、実施例1より実施例2、3の方が大きい。その理由は、実施例1においては、熱間鍛造の途中、即ち断面が45mm角になったところで、昇温加熱して1200℃で、5分間の再加熱保持をしたため、この時点で全く加工していない組織レベルに組織が戻ってしまったからである。これは、別途、フォーマスター試験により、完全にα→γ変態を終了してしまうことを確認していることからもいえる。この再加熱後での熱間鍛造により導入された塑性相当ひずみの量e(断面:45mm角から仕上がりの断面38mm角までの減面率は、28.7%である)は、e=0.338と算出されるのに対して、熱間鍛造開始後、途中で再加熱なしで断面:95mm角から38mm角まで熱間鍛造を継続した実施例2及び実施例3においては、導入された塑性相当ひずみの量eは、e=1.83と算出される。このとき、実施例2、3は、実施例1よりも若干引張強度が高く、一方全伸びは若干低くなっている。 Next, looking at the amount of plastic equivalent strain to the finished steel material by hot forging, Examples 2 and 3 are larger than Example 1. The reason is that in Example 1, during the hot forging, that is, when the cross section became 45 mm square, the temperature was raised and heated and reheated at 1200 ° C. for 5 minutes. This is because the organization has returned to a non-organizational level. This can also be said from the fact that it has been confirmed that the α → γ transformation is completely completed by a formaster test. The amount of plastic equivalent strain e introduced by hot forging after this reheating (cross section: the area reduction rate from the 45 mm square to the finished 38 mm square is 28.7%) is e = 0. In contrast to the calculation of 338, in Example 2 and Example 3 in which hot forging was continued from 95 mm square to 38 mm square without reheating in the middle after the start of hot forging, the introduced plasticity The amount of equivalent strain e is calculated as e = 1.83. At this time, Examples 2 and 3 have slightly higher tensile strength than Example 1, while total elongation is slightly lower.
なお、以上の実施例及び比較例の試験は、試験材のサイズが実験室レベルの規模のものであるが、その他の条件は全て、現状の実機操業に適用可能な技術条件に匹敵することは、明らかである。 The tests in the above examples and comparative examples are of a test material size on the laboratory level, but all other conditions are comparable to the technical conditions applicable to the current actual machine operation. ,it is obvious.
これに対して、比較例の試験結果は以下の通りである。 On the other hand, the test results of the comparative example are as follows.
比較例1は、顕微鏡組織に関しては主相は残留オーステナイトであるが、第2相がフェライトとなっており、また前記図7に示したSEMによる組織写真からも判るように、ラスの周囲に球状化セメンタイトが生成している。このように、本願発明の鋼材の範囲外にある。比較例1は、全伸び(El)は優れているが、引張強度(TS)が933MPaと低く、合否判定基準の1300MPaを満たさず、上部棚エネルギー(USE)も31J/cm2と低く、合否判定基準の100J/cm2を大きく下回っている。比較例1の引張強度がこのように著しく低い原因は、熱間鍛造(実施例1は熱間鍛造のままで、非調質である)後の焼鈍処理により、残留オーステナイト中のCが球状化セメンタイトとして排出されたために、主相の残留オーステナイトの硬さが著しく低下するとともに、第2相がフェライトとなり、ラスマルテンサイトが消失したからである。 In Comparative Example 1, the main phase is retained austenite with respect to the microstructure, but the second phase is ferrite, and as can be seen from the structure photograph by SEM shown in FIG. Cementite is formed. Thus, it is outside the range of the steel material of the present invention. In Comparative Example 1, the total elongation (El) is excellent, but the tensile strength (TS) is as low as 933 MPa, the acceptance criteria of 1300 MPa is not satisfied, and the upper shelf energy (USE) is also as low as 31 J / cm 2. This is far below the judgment standard of 100 J / cm 2 . The reason why the tensile strength of Comparative Example 1 is so low is that C in the retained austenite is spheroidized by the annealing treatment after hot forging (Example 1 is still hot forged and is not tempered). This is because the hardness of the retained austenite of the main phase is remarkably lowered and the second phase becomes ferrite and the lath martensite disappears because it is discharged as cementite.
また、ビッカース硬さの10kg試験結果においても、実施例1(HV=416)と比較例1(HV=299)との差が明確に現れている。
Also in Vickers hardness 10kg test results, the difference between Example 1 (HV = 416) Comparative Example 1 (HV = 299) are to appear clearly.
耐衝撃エネルギー吸収特性については、比較例1は熱間鍛造後に、焼鈍処理をしているので、従来技術の知見によれば実施例1の熱間鍛造ままよりも優れていると予測されるところであるが、逆に実施例1の方が耐衝撃エネルギー吸収特性が優れている。これは、実施例1では主相が、ラスマルテンサイトであって、一部、体積分率5%以下の残留オーステナイトを含む組織で短径及び長径共に微細となっているためであると考えられる。
As for the impact energy absorption characteristics, since Comparative Example 1 is annealed after hot forging, it is predicted that it is superior to the hot forging of Example 1 according to the knowledge of the prior art. On the contrary, the impact energy absorption characteristics of Example 1 are superior. This is presumably because, in Example 1, the main phase is lath martensite, and a part of the structure containing residual austenite with a volume fraction of 5% or less is fine in both the minor axis and the major axis. .
比較例2は、顕微鏡組織に関して主相はフェライトであり、第2相が残留オーステナイトとなっている。 In Comparative Example 2, the main phase of the microstructure is ferrite, and the second phase is retained austenite.
比較例2のSEMによる組織写真を前記図8に示した。比較例2は、全伸び(El)は優れているが、引張強度(TS)が1142MPaと低く、合否判定基準の1300MPaを下回っている。比較例2の引張強度がこのように実施例1〜3に比べて低く、また、上記金属組織を示す理由は、熱間鍛造の後に、更に温間溝ロール圧延並びに焼鈍及び過時効処理を施しているためであると考えられる。 The structure photograph by the SEM of the comparative example 2 was shown in the said FIG. In Comparative Example 2, the total elongation (El) is excellent, but the tensile strength (TS) is as low as 1142 MPa, which is below the pass / fail criterion 1300 MPa. The reason why the tensile strength of Comparative Example 2 is thus lower than that of Examples 1 to 3 and the above metal structure is shown is that after hot forging, further hot groove rolling, annealing and overaging treatment are performed. It is thought that this is because.
冷却速度を変えて生成したラスマルテンサイトにおけるEBSD粒界マップを図14に、それから計算により得られた各傾角の粒界面積当たり長さを図15に示す。これから、本発明材の特徴として
大傾角(15°−180°)粒界面積当たり長さ>0.6μm/μm2
中傾角(5°−15°)粒界面積当たり長さ>0.2μm/μm2
小傾角(1.5°−5°)粒界面積当たり長さ>0.3μm/μm2
であることが認められる。
FIG. 14 shows an EBSD grain boundary map in lath martensite generated by changing the cooling rate, and FIG. 15 shows the length per grain interface area of each inclination obtained from the calculation. From this, as a feature of the material of the present invention, the length per large interfacial area (15 ° -180 °) grain interface area> 0.6 μm / μm 2
Medium tilt angle (5 ° -15 °) Length per grain interfacial area> 0.2 μm / μm 2
Small tilt angle (1.5 ° -5 °) Length per grain interface area> 0.3 μm / μm 2
It is recognized that
これは単に1200℃のγ域から種々の冷却速度で冷却するのみで、マルテンサイトラス構造が形成され、それを構成している粒界が高密度の大角粒界からなっている特徴を示すものである。このことが、高強度、高延性、高靭性の要因であると考えられる。 This is a characteristic that a martensitic structure is formed only by cooling at a γ region of 1200 ° C. at various cooling rates, and the grain boundaries constituting it are composed of high-density, large-angle grain boundaries. is there. This is considered to be a factor of high strength, high ductility, and high toughness.
本発明により、建造物や橋梁等の構造物、自動車の足回り鋼材、機械用歯車等部品に使用される鋼材であって、高強度且つ高延性で、衝撃エネルギー吸収能に優れた厚鋼板や棒鋼・鋼線等の鋼材を、省資源、省エネルギーかつ、低コスト、省工程で製造することができる。 According to the present invention, steel materials used for structures such as buildings and bridges, undercarriage steel materials for automobiles, mechanical gears, etc., which are steel plates having high strength and high ductility and excellent shock energy absorption capability, Steel materials such as steel bars and steel wires can be manufactured in resource-saving, energy-saving, low-cost, and process-saving processes.
Claims (4)
組織は、主相がラスマルテンサイトであって、一部に体積分率5%以下の残留オーステナイトからなることを特徴とする強度・延性・靭性に優れた高強度非調質鋼材。 Chemical composition is mass%, C: 0.05-0.20%, Si: 1.0-2.2%, Mn: 3.0-5.5%, Al: 0.0005-0.08 %, Non-tempered steel with the balance being Fe and inevitable impurities,
Organization, main phase a lath martensite, high strength non-heat treated steel with excellent strength and ductility and toughness, characterized in that it consists of a volume fraction of 5% or less of retained austenite in a part.
It is manufactured by heat-maintaining within the range of 1000-1300 degreeC, and performing the hot work whose plastic equivalent strain e is 0.34 or more to this, It is any one of Claim 1 to 3 characterized by the above-mentioned. method for producing a non-heat treated steel excellent in strength, ductility and toughness of.
Priority Applications (1)
Application Number | Priority Date | Filing Date | Title |
---|---|---|---|
JP2010249102A JP6021094B2 (en) | 2010-11-05 | 2010-11-05 | High-strength non-heat treated steel material excellent in strength, ductility and toughness and method for producing the same |
Applications Claiming Priority (1)
Application Number | Priority Date | Filing Date | Title |
---|---|---|---|
JP2010249102A JP6021094B2 (en) | 2010-11-05 | 2010-11-05 | High-strength non-heat treated steel material excellent in strength, ductility and toughness and method for producing the same |
Publications (2)
Publication Number | Publication Date |
---|---|
JP2012102346A JP2012102346A (en) | 2012-05-31 |
JP6021094B2 true JP6021094B2 (en) | 2016-11-02 |
Family
ID=46393084
Family Applications (1)
Application Number | Title | Priority Date | Filing Date |
---|---|---|---|
JP2010249102A Expired - Fee Related JP6021094B2 (en) | 2010-11-05 | 2010-11-05 | High-strength non-heat treated steel material excellent in strength, ductility and toughness and method for producing the same |
Country Status (1)
Country | Link |
---|---|
JP (1) | JP6021094B2 (en) |
Families Citing this family (5)
Publication number | Priority date | Publication date | Assignee | Title |
---|---|---|---|---|
JP5747249B2 (en) * | 2011-04-15 | 2015-07-08 | 国立研究開発法人物質・材料研究機構 | High-strength steel material excellent in strength, ductility and energy absorption capacity and its manufacturing method |
JP5747250B2 (en) * | 2011-04-25 | 2015-07-08 | 国立研究開発法人物質・材料研究機構 | High-strength steel material excellent in strength, ductility and impact energy absorption ability, and method for producing the same |
JP6327737B2 (en) | 2013-07-09 | 2018-05-23 | 国立研究開発法人物質・材料研究機構 | Martensitic steel and manufacturing method thereof |
CN104846274B (en) * | 2015-02-16 | 2017-07-28 | 重庆哈工易成形钢铁科技有限公司 | Hot press-formed use steel plate, hot press-formed technique and hot press-formed component |
CN114855098A (en) * | 2022-04-22 | 2022-08-05 | 东北大学 | High-strength medium manganese steel for engineering machinery and preparation method thereof |
-
2010
- 2010-11-05 JP JP2010249102A patent/JP6021094B2/en not_active Expired - Fee Related
Also Published As
Publication number | Publication date |
---|---|
JP2012102346A (en) | 2012-05-31 |
Similar Documents
Publication | Publication Date | Title |
---|---|---|
JP6465266B1 (en) | Hot rolled steel sheet and manufacturing method thereof | |
JP6849536B2 (en) | High-strength steel sheet and its manufacturing method | |
JP5176885B2 (en) | Steel material and manufacturing method thereof | |
JP6468408B2 (en) | H-section steel and its manufacturing method | |
RU2502820C1 (en) | Plate steel characterised by low ratio between yield point and ultimate strength, high strength and high uniform relative elongation, and method for its manufacture | |
JP6103165B1 (en) | Hot press-formed parts | |
JP4324225B1 (en) | High strength cold-rolled steel sheet with excellent stretch flangeability | |
JP5126399B2 (en) | High-strength cold-rolled steel sheet with excellent stretch flangeability and manufacturing method thereof | |
JP5803836B2 (en) | Hot pressed steel plate member, its manufacturing method and hot pressed steel plate | |
JP5439973B2 (en) | High-strength thick steel plate having excellent productivity and weldability and excellent drop weight characteristics after PWHT, and method for producing the same | |
MX2014010648A (en) | High-strength cold-rolled steel sheet and process for manufacturing same. | |
CN105008570A (en) | Thick, tough, high tensile strength steel plate and production method therefor | |
EP3222743B1 (en) | Rolled steel bar or rolled wire material for cold-forged component | |
JP6811694B2 (en) | Steel plate and its manufacturing method | |
JP6809524B2 (en) | Ultra-low yield ratio high-strength thick steel sheet and its manufacturing method | |
JP4983082B2 (en) | High-strength steel and manufacturing method thereof | |
JP6988836B2 (en) | Ultra-low yield ratio high-strength thick steel sheet and its manufacturing method | |
JP6795048B2 (en) | Non-treated low yield ratio high-strength thick steel sheet and its manufacturing method | |
JP4207334B2 (en) | High-strength steel sheet with excellent weldability and stress corrosion cracking resistance and method for producing the same | |
JP5786720B2 (en) | High tensile thick steel plate having a tensile strength of 780 MPa or more and method for producing the same | |
JP2009173959A (en) | High-strength steel sheet and producing method therefor | |
JP6021094B2 (en) | High-strength non-heat treated steel material excellent in strength, ductility and toughness and method for producing the same | |
KR20130046941A (en) | High strength steel sheet and method of manufacturing the steel sheet | |
US20160333448A1 (en) | Steel and method of manufacturing the same | |
JP5477089B2 (en) | Manufacturing method of high strength and high toughness steel |
Legal Events
Date | Code | Title | Description |
---|---|---|---|
A621 | Written request for application examination |
Free format text: JAPANESE INTERMEDIATE CODE: A621 Effective date: 20131105 |
|
A711 | Notification of change in applicant |
Free format text: JAPANESE INTERMEDIATE CODE: A711 Effective date: 20131105 |
|
A521 | Request for written amendment filed |
Free format text: JAPANESE INTERMEDIATE CODE: A821 Effective date: 20131112 |
|
A521 | Request for written amendment filed |
Free format text: JAPANESE INTERMEDIATE CODE: A523 Effective date: 20140114 |
|
A977 | Report on retrieval |
Free format text: JAPANESE INTERMEDIATE CODE: A971007 Effective date: 20141020 |
|
A131 | Notification of reasons for refusal |
Free format text: JAPANESE INTERMEDIATE CODE: A131 Effective date: 20141104 |
|
RD04 | Notification of resignation of power of attorney |
Free format text: JAPANESE INTERMEDIATE CODE: A7424 Effective date: 20141127 |
|
A521 | Request for written amendment filed |
Free format text: JAPANESE INTERMEDIATE CODE: A523 Effective date: 20141222 |
|
A02 | Decision of refusal |
Free format text: JAPANESE INTERMEDIATE CODE: A02 Effective date: 20150303 |
|
A521 | Request for written amendment filed |
Free format text: JAPANESE INTERMEDIATE CODE: A523 Effective date: 20150508 |
|
A911 | Transfer to examiner for re-examination before appeal (zenchi) |
Free format text: JAPANESE INTERMEDIATE CODE: A911 Effective date: 20150518 |
|
A912 | Re-examination (zenchi) completed and case transferred to appeal board |
Free format text: JAPANESE INTERMEDIATE CODE: A912 Effective date: 20150710 |
|
A521 | Request for written amendment filed |
Free format text: JAPANESE INTERMEDIATE CODE: A523 Effective date: 20160720 |
|
A61 | First payment of annual fees (during grant procedure) |
Free format text: JAPANESE INTERMEDIATE CODE: A61 Effective date: 20160927 |
|
R150 | Certificate of patent or registration of utility model |
Ref document number: 6021094 Country of ref document: JP Free format text: JAPANESE INTERMEDIATE CODE: R150 |
|
R250 | Receipt of annual fees |
Free format text: JAPANESE INTERMEDIATE CODE: R250 |
|
R250 | Receipt of annual fees |
Free format text: JAPANESE INTERMEDIATE CODE: R250 |
|
R250 | Receipt of annual fees |
Free format text: JAPANESE INTERMEDIATE CODE: R250 |
|
LAPS | Cancellation because of no payment of annual fees |