JPH0518897B2 - - Google Patents
Info
- Publication number
- JPH0518897B2 JPH0518897B2 JP10229888A JP10229888A JPH0518897B2 JP H0518897 B2 JPH0518897 B2 JP H0518897B2 JP 10229888 A JP10229888 A JP 10229888A JP 10229888 A JP10229888 A JP 10229888A JP H0518897 B2 JPH0518897 B2 JP H0518897B2
- Authority
- JP
- Japan
- Prior art keywords
- less
- tio
- steel
- ductility
- ferrite
- 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 - Lifetime
Links
- 229910010413 TiO 2 Inorganic materials 0.000 claims description 34
- 229910000831 Steel Inorganic materials 0.000 claims description 28
- 239000010959 steel Substances 0.000 claims description 28
- 239000002245 particle Substances 0.000 claims description 18
- 229910000859 α-Fe Inorganic materials 0.000 claims description 17
- GWEVSGVZZGPLCZ-UHFFFAOYSA-N Titan oxide Chemical compound O=[Ti]=O GWEVSGVZZGPLCZ-UHFFFAOYSA-N 0.000 claims description 14
- 239000002131 composite material Substances 0.000 claims description 13
- 229910000734 martensite Inorganic materials 0.000 claims description 9
- 239000011159 matrix material Substances 0.000 claims description 6
- 229910052750 molybdenum Inorganic materials 0.000 claims description 6
- 239000006185 dispersion Substances 0.000 claims description 5
- 229910052804 chromium Inorganic materials 0.000 claims description 4
- 239000012535 impurity Substances 0.000 claims description 4
- 229910052748 manganese Inorganic materials 0.000 claims description 4
- 229910052742 iron Inorganic materials 0.000 claims description 2
- 230000009977 dual effect Effects 0.000 claims 1
- 239000002436 steel type Substances 0.000 description 14
- 239000010936 titanium Substances 0.000 description 10
- 229910045601 alloy Inorganic materials 0.000 description 9
- 239000000956 alloy Substances 0.000 description 9
- 239000000463 material Substances 0.000 description 8
- 239000000843 powder Substances 0.000 description 8
- 238000000137 annealing Methods 0.000 description 7
- 239000000203 mixture Substances 0.000 description 6
- 239000000446 fuel Substances 0.000 description 5
- 239000008358 core component Substances 0.000 description 4
- 230000007423 decrease Effects 0.000 description 4
- 230000000694 effects Effects 0.000 description 4
- 238000004519 manufacturing process Methods 0.000 description 4
- 238000000034 method Methods 0.000 description 4
- 239000006104 solid solution Substances 0.000 description 4
- 238000005482 strain hardening Methods 0.000 description 4
- 238000005728 strengthening Methods 0.000 description 4
- 229910052721 tungsten Inorganic materials 0.000 description 4
- 239000000654 additive Substances 0.000 description 3
- 230000000052 comparative effect Effects 0.000 description 3
- 239000000306 component Substances 0.000 description 3
- 230000007774 longterm Effects 0.000 description 3
- 238000009864 tensile test Methods 0.000 description 3
- XKRFYHLGVUSROY-UHFFFAOYSA-N Argon Chemical compound [Ar] XKRFYHLGVUSROY-UHFFFAOYSA-N 0.000 description 2
- DGAQECJNVWCQMB-PUAWFVPOSA-M Ilexoside XXIX Chemical compound C[C@@H]1CC[C@@]2(CC[C@@]3(C(=CC[C@H]4[C@]3(CC[C@@H]5[C@@]4(CC[C@@H](C5(C)C)OS(=O)(=O)[O-])C)C)[C@@H]2[C@]1(C)O)C)C(=O)O[C@H]6[C@@H]([C@H]([C@@H]([C@H](O6)CO)O)O)O.[Na+] DGAQECJNVWCQMB-PUAWFVPOSA-M 0.000 description 2
- 229910004298 SiO 2 Inorganic materials 0.000 description 2
- 230000000996 additive effect Effects 0.000 description 2
- 239000003795 chemical substances by application Substances 0.000 description 2
- 238000005253 cladding Methods 0.000 description 2
- 238000002156 mixing Methods 0.000 description 2
- 238000001556 precipitation Methods 0.000 description 2
- 239000000047 product Substances 0.000 description 2
- 229910052708 sodium Inorganic materials 0.000 description 2
- 239000011734 sodium Substances 0.000 description 2
- 239000000126 substance Substances 0.000 description 2
- 230000008961 swelling Effects 0.000 description 2
- 238000005496 tempering Methods 0.000 description 2
- RTAQQCXQSZGOHL-UHFFFAOYSA-N Titanium Chemical compound [Ti] RTAQQCXQSZGOHL-UHFFFAOYSA-N 0.000 description 1
- 238000005275 alloying Methods 0.000 description 1
- 229910052786 argon Inorganic materials 0.000 description 1
- 230000000712 assembly Effects 0.000 description 1
- 238000000429 assembly Methods 0.000 description 1
- 229910000963 austenitic stainless steel Inorganic materials 0.000 description 1
- 238000006243 chemical reaction Methods 0.000 description 1
- 238000005097 cold rolling Methods 0.000 description 1
- 238000001816 cooling Methods 0.000 description 1
- 238000005260 corrosion Methods 0.000 description 1
- 230000007797 corrosion Effects 0.000 description 1
- 238000005261 decarburization Methods 0.000 description 1
- 230000003247 decreasing effect Effects 0.000 description 1
- 230000003009 desulfurizing effect Effects 0.000 description 1
- 238000001125 extrusion Methods 0.000 description 1
- 238000005242 forging Methods 0.000 description 1
- 239000007789 gas Substances 0.000 description 1
- 238000010438 heat treatment Methods 0.000 description 1
- 238000001192 hot extrusion Methods 0.000 description 1
- 238000009863 impact test Methods 0.000 description 1
- 230000003993 interaction Effects 0.000 description 1
- 229910000765 intermetallic Inorganic materials 0.000 description 1
- 229910001338 liquidmetal Inorganic materials 0.000 description 1
- 150000001247 metal acetylides Chemical class 0.000 description 1
- 230000003647 oxidation Effects 0.000 description 1
- 238000007254 oxidation reaction Methods 0.000 description 1
- 229910001175 oxide dispersion-strengthened alloy Inorganic materials 0.000 description 1
- 239000002244 precipitate Substances 0.000 description 1
- 239000002994 raw material Substances 0.000 description 1
- 229920006395 saturated elastomer Polymers 0.000 description 1
- 238000001179 sorption measurement Methods 0.000 description 1
- 229910001220 stainless steel Inorganic materials 0.000 description 1
- 238000003756 stirring Methods 0.000 description 1
- 229910052719 titanium Inorganic materials 0.000 description 1
- OGIDPMRJRNCKJF-UHFFFAOYSA-N titanium oxide Inorganic materials [Ti]=O OGIDPMRJRNCKJF-UHFFFAOYSA-N 0.000 description 1
- 230000007704 transition Effects 0.000 description 1
Landscapes
- Powder Metallurgy (AREA)
- Treatment Of Steel In Its Molten State (AREA)
Description
＜産業上の利用分野＞
本発明は、優れた高温強度を有し、しかも延性
および靱性に優れた原子炉用分散強化フエライト
鋼に関するものである。
本発明の分散強化フエライト鋼は、原子炉、特
に高速増殖炉の炉心で使用される炉心構成要素
（例えば、燃料集合体、制御棒、反射体等）や高
速炉の機器構造物（例えば、機器容器部材、冷却
系配管部材）などに好ましく利用できる。
＜従来技術＞
原子炉、特に高速増殖炉の炉心構成部材に用い
られる材料には、高温強度、ナトリウムとの共存
在、耐中性子照射特性、加工性、溶接性、燃料と
の相互作用など様々な特性が要求されるが、特に
高温強度と耐中性子照射特性がその使用寿命を決
定する上で重要である。
従来より炉心部材としてはSUS304や316など
のオーステナイトステンレス鋼が用いられてきた
が、耐スエリング性や照射クリープ特性など高速
中性子に対する耐久性に限界があり、燃料の長寿
命化を達成するには適していないことが明らかに
なつている。
一方、フエライト鋼はオーステナイトステンレ
ス鋼に比べ、格段に優れた耐照射特性を有するも
のの、高温強度が低い欠点がある。高温強度向上
の方法の一つとして、微細酸化物粒子による分散
強化法が古くより知られている。この方法を用い
たフエライト鋼としては「液体金属高速中性子増
殖炉用の分散強化フエライト型合金」（特公昭60
−8296号、以下先行特許という）があり、高温ク
リープ破断強度と耐中性子照射特性に優れたもの
が得られている。
＜発明が解決しようとする課題＞
しかし、上記先行特許による合金は強度が高い
反面、延性が低く、また延性・脆性遷移温度が20
℃付近と高いため、即ち室温での衝撃値が著しく
低いため、わずか10数％程度の冷間圧延で割れが
生ずる。その結果、高速炉の炉心部材、例えば燃
料被覆管あるいはラツパー管などの高い寸法精度
が要求される薄肉管の経済的な製管は、上記先行
特許による合金では困難である。さらに高速炉の
使用温度350〜700℃において亀裂が極めて伝播し
やすい低延性材料であり、分散強化材本来の特性
が生かされていない。
そこで本発明は、高速増殖炉の炉心部材として
要求される性質、すなわち高温強度や耐スエリン
グ性に優れ、しかも延性、靱性、製管性なども良
好な分散強化フエライト鋼を提供することを目的
としてなされたものである。
＜課題を解決するための手段＞
すなわち、本発明の延性と靱性に優れた原子炉
用分散強化フエライト鋼の第１の実施態様は、重
量％でＣ：0.1％以下、Si：0.1％以下、Mn：0.1
％以下、Cr：12〜20％、Mo＋Ｗ＝0.1〜4.4％、
Ｏ（Y2O3およびTiO2分は除く）：0.01％以下、残
部がFeおよび不可避不純物からなり、かつ平均
粒径1000Å以下のY2O3とTiO2による複合酸化物
粒子がY2O3＋TiO2：0.1〜1.0％、分子比で
TiO2／Y2O3＝0.5〜2.0の範囲で基地（マトリツ
クス）に均一に分散されているフエライト単相あ
るいはフエライト／マルテンサイト２相組織であ
ることを特徴とするものである。
さらに本発明の延性と靱性に優れた原子炉用分
散強化フエライト鋼の第２の実施態様は、重量％
でＣ：0.05％〜0.25％、Si：0.1％以下、Mn：0.1
％以下、Cr：８〜12％（但し、12％は含まず）、
Mo＋Ｗ：0.1〜4.0％、Ｏ（Y2O3およびTiO2分は
除く）：0.01％以下、残部がFeおよび不可避不純
物からなり、かつ平均粒径1000Å以下のY2O3と
TiO2による複合酸化物粒子がY2O3＋TiO2＝0.1
〜1.0％、分子比でTiO2／Y2O3＝0.5〜2.0の範囲
で基地に均一に分散されている焼戻しマルテンサ
イト単相組織であることを特徴とするものであ
る。
以下、本発明の分散強化フエライト鋼の化学成
分およびその限定理由について述べる。このうち
Y2O3とTiO2の複合添加が本発明の最重要ポイン
トである。
Y2O3は、基地に均一分散されることによりク
リープ破断強度を向上させる効果を有する最も重
要な成分である。しかしながら、Y2O3単独では、
基地に固溶する少量のSiやMn等と容易に複合酸
化物を形成して粗大化しやすくなる。また、Y2
O3そのものは基地との整合性が悪く、Y2O3を多
量に添加してもクリープ破断強度は向上せず、む
しろ延性や靱性が低下する。
Y2O3とTiO2の安定な複合添加物
Y2O3・TiO2を形成させて初めて高い強度が得
られる。
複合添加物Y2O3・TiO2即ちY2TiO5は、基地
合金組成粉末とY2O3微粉末を機械的に混合する
過程においてTiO2微粉末を添加することにより
形成される。Y2O3単独よりもY2O3・TiO2のほう
がエネルギー上安定であるため、混合時にすべて
のY2O3とTiO2が反応する。また、あらかじめ調
製されたY2O3・TiO2複合酸化物を使用すること
も可能である。
前述した先行特許においても「イツトリア
（Y2O3）組成物中の他の成分たとえばチタンと結
合してY2Ti2O7のような相を形成することができ
る。」と記載されている。しかし、この方法では、
基地中に固溶したTiとY2O3粒子が反応して複合
酸化物ができるため、複合酸化物の組成が不均一
となり、Ti濃度の高いものや不足したものが生
じる。このような酸化物は、かえつて熱的に不安
定であるため製管時に高温軟化処理を施すと凝集
粗大化し、クリープ破断強度を低下させる。ま
た、Y2O3と反応しない過剰なTiは、TiO2酸化物
単体として析出する。TiO2は高温使用中に粗大
化しやすく延性低下の原因となる。また、一度粗
大化してしまうと、高温軟化焼鈍を施しても合金
の延性は回復しない。
本発明では、Y2O3とTiO2を分子比で0.5〜2.0
の範囲で反応させることにより、安定でしかも基
地との整合性の良い複合酸化物を均一に分散させ
ることができる。
また、添加した全てのTiO2とY2O3との反応に
より安定な複合酸化物となつているので、高温の
軟化焼鈍により延性は加工前のそれに回復する。
（Y2O3＋TiO2）量は、高温強度を向上させるの
に、最低0.1％以上必要である。一方、（Y2O3＋
TiO2）の添加量を多くすれば、クリープ破断強
度は高くなるが、その効果は、１％で飽和するの
で添加量の上限を１％とする。
また、酸化物粒径については、粒径が1000Åを
越えるとクリープ強度を高める効果が著しく低下
するので1000Å以下に限定する。
Ｃは、組織の安定性を左右する重要な元素であ
り、要求特性によつてその成分範囲は異なる。
延性を重視した場合は、組織をフエライト単相
か少量の焼戻しマルテンサイトを含むフエライ
ト／マルテンサイト２相にする必要があり、Cr
含有量が12〜20％の第１の実施態様では、Ｃ量は
0.1％以下、好ましくは0.01％未満に抑える必要
がある。また、Ｃ含有量を低減することにより、
基地に固溶するMo、Ｗの固溶量を多くすること
ができ、長時間側のクリープ強度を高めることが
可能となる。
一方、靱性を重視した場合は、組織を安定な焼
戻しマルテンサイト単相にする必要があり、Cr
含有量が８〜12％の第２の実施態様では、Ｃ量の
下限は0.05％となる。この焼戻しマルテンサイト
組織は、1000〜1150℃の焼ならし＋700〜800℃の
焼戻し処理により得られる。Ｃ含有量が多くなる
ほど炭化物（M23C6、M6Ｃなど）の析出量が多
くなり高温強度が高くなるが、0.25％より多量に
添加すると加工性が悪くなる。よつてこの場合の
Ｃ含有量は0.05〜0.25％に限定する。
Crは、Ｃ量とのバランスからＣ含有量が0.05％
未満の場合、12〜20％の範囲でフエライト相、Ｃ
含有量が0.05〜0.25％の場合は、８〜12％の範囲
で焼戻しマルテンサイト相を安定させることがで
きる。また、８％よりも少ないと高温（600〜700
℃）でのナトリウム中脱炭抵抗性および耐食性が
悪くなる。一方、20％以上では、靱性、延性が低
下するので上限を20％とする。
Siは、脱酸剤として必要な元素であるが、使用
中にY2O3粒子と反応してY2O3とSiO2の複合酸化
物を形成しやすい。この複合酸化物は、粗大化速
度が大きいため、クリープ破断強度を低下させ
る。また低Si化によつて製品の表面性状を良好に
し、SiO2介在物量を少なくすることもできるの
で、0.1％以下に抑える。
Mnは、脱酸・脱硫剤として働き、熱間加工性
の改善にも有効な元素であるが、多量に添加する
とSi同様粗大化しやすいY2O3との複合酸化物を
形成するので、0.1％以下に限定する。
MoとＷは合金中に固溶し、高温強度を向上さ
せる重要な元素であり、総量で0.1％以上添加す
る必要がある。MoとＷ量を多くすれば、固溶強
化作用、炭化物析出強化作用（M23C6、M6Ｃな
ど）、金属間化合物析出強化作用により、クリー
プ破断強度が向上するが、Mo＋Ｗで4.0％を超え
るとδフエライト量が多くなり、かえつて強度も
低下するので、4.0％を上限とする。
特に高い強度を得るには、Mo、Ｗの複合によ
りMo当量（Mo＋1/2Ｗ）が1.2〜1.6％となる組
合わせが良い。
Ｏは、原料粉末上への吸着あるいは酸化により
必然的に少量含まれる元素であるが、0.01％を超
えると靱性が著しく低下する。また、少量のSiや
Mnと介在物を形成しやすくなるので、その上限
を0.01％とする。
＜実施例＞
以下に本発明について実施例を挙げて説明す
る。
表１に供試材の化学成分を示す。
表１で、鋼種No.１〜３は本発明鋼（）（本発
明の第１の実施態様に該当）、鋼種No.４〜６は本
発明鋼（）（本発明の第２の実施態様に該当）、
鋼種No.７〜13は本発明鋼（）の比較鋼、鋼種14
〜20は本発明鋼（）の比較鋼である。比較鋼に
はそれぞれ重要な添加物であるTiO2、Y2O3、
（Mo＋Ｗ）、Ｃ量が本発明の範囲とはずれている
もの、あるいはTiO2の代わりにTiを添加したも
のを用いた。このうちNo.13は先行特許で提案され
た合金に相当する。
各鋼とも平均粒径1μm以下の元素粉あるいは合
金粉と平均粒径1000Å以下の酸化物粉末を目的組
成に調合し、高エネルギーアトライター中に装入
後、高純度アルゴンガス雰囲気中で攪拌して機械
的に合金化を行つた。アトライターの回転数は
200〜300rpm、攪拌時間は24〜48hrである。得ら
れた合金化粉末を空気にさらすことなくSUS製
の筒状容器に真空封入し、900〜1200℃で８〜
15：１の押出比で熱間押出した。
各熱押棒材を10mm厚の板材に鍛造した後、950
〜1200℃で焼ならしを行い、鋼種No.１〜３及び７
〜13については焼ならしたままのもの、鋼種No.４
〜６および14〜20については焼ならし後、750〜
820℃の焼もどし熱処理を施したものを供試材と
した。
これらの供試材から2t×6W×30GL（mm）の板
状引張試験片を採取し、650℃クリープ破断試験
及び常温引張試験を行つた。また、5t×10W×55
mm（２mmＶノツチ）のシヤルピー衝撃試験片を
採取し、衝撃特性を調べた。さらに、10（mm）厚
板材を20％冷間圧延後1200℃×1hr焼鈍したもの
より2t×6W×30GLmmの板状試験片を採取し、常
温引張試験を行い、焼鈍後の引張延性の変化を調
べた。
それらの試験結果より、650℃×103hrでのクリ
ープ破断応力、常温引張伸び、20℃でのシヤルピ
ー衝撃値及び冷間加工＋焼鈍後の引張伸びを表２
にまとめて示す。
本発明鋼（），（）ともに、比較鋼に比べ
650℃クリープ破断強度、常温での延性、20℃で
のシヤルピー衝撃特性、冷間加工＋焼鈍後の延性
に優れていることがわかる。特に本発明鋼（）
は延性、本発明鋼（）は靱性に優れている。比
較鋼のうち鋼種No.７とNo.14は、TiO2添加量が不
足しているため、分散酸化物粒子が不安定で強度
が低い。
鋼種No.８とNo.15は、強度は比較的高いものの、
TiO2添加量が過剰であるため、分散酸化物粒子
の凝集、粗大化が起こり延性、靱性が低くなつて
いる。
鋼種No.９とNo.16は、TiO2無添加、鋼種No.10と
17はY2O3無添加でいずれの場合も強度低下が著
しい。このことは、Y2O3、TiO2単独添加では分
散物が不安定であり、Y2O3とTiO2共存下で初め
て分散物が安定化され強度が高くなることを示し
ている。
鋼種No.11と18は、TiO2の代わりにTiを添加し
たものである。いずれも強度は比較的高いもの
の、延性特に冷間加工＋焼鈍後の延性が低い。こ
れは、TiはY2O3と反応して複合酸化物を作る以
外に、粗大化しやすいTiO2単体として析出する
ため、TiO2粒子が応力集中源となり延性や靱性
が低下することを示している。
鋼種No.12と19は、（Mo＋Ｗ）含有量が不足し
ているため、強度が低く表２には示していないが
特に長時間側の強度低下が大きい。また、固溶元
素が少ないため、基地の強化がなされず、他の鋼
種に比べ靱性が低い。
鋼種No.13は先行特許で提案された合金である
が、鋼種No.11，18と同様TiO2の代わりにTiを添
加しているため、本発明鋼に比べ、延性や靱性が
低い。
鋼種No.20は、本発明鋼（）の範囲よりＣ含有
量が低い。このためフエライト相が多く現れるた
め、延性や靱性が著しく低くなつている。
<Industrial Application Field> The present invention relates to a dispersion-strengthened ferrite steel for nuclear reactors, which has excellent high-temperature strength, as well as excellent ductility and toughness. The dispersion-strengthened ferritic steel of the present invention can be used in core components (e.g., fuel assemblies, control rods, reflectors, etc.) used in the core of nuclear reactors, particularly fast breeder reactors, and equipment structures of fast reactors (e.g., equipment It can be preferably used for container members, cooling system piping members), etc. <Prior art> Materials used for core components of nuclear reactors, especially fast breeder reactors, have various characteristics such as high temperature strength, coexistence with sodium, neutron irradiation resistance, workability, weldability, and interaction with fuel. Among the required properties, high-temperature strength and neutron irradiation resistance are particularly important in determining its service life. Traditionally, austenitic stainless steels such as SUS304 and 316 have been used for core components, but they have limited durability against fast neutrons such as swelling resistance and irradiation creep properties, making them unsuitable for achieving longer fuel life. It is clear that this is not the case. On the other hand, although ferritic steel has much better irradiation resistance than austenitic stainless steel, it has the disadvantage of low high-temperature strength. Dispersion strengthening using fine oxide particles has long been known as one of the methods for improving high-temperature strength. The ferritic steel produced using this method is the ``dispersion-strengthened ferrite-type alloy for liquid metal fast neutron breeder reactors'' (Special Publication Publication in 1983).
No. 8296 (hereinafter referred to as the prior patent), a product with excellent high-temperature creep rupture strength and neutron irradiation resistance has been obtained. <Problem to be solved by the invention> However, although the alloy according to the prior patent mentioned above has high strength, it has low ductility, and the ductile-brittle transition temperature is 20
Since the impact value is extremely low at around ℃, that is, the impact value at room temperature is extremely low, cracks occur after only about 10% cold rolling. As a result, it is difficult to economically manufacture thin-walled tubes that require high dimensional accuracy, such as core members of fast reactors, such as fuel cladding tubes or Lapper tubes, using the alloys according to the prior patents mentioned above. Furthermore, it is a low ductility material that is extremely susceptible to crack propagation at the operating temperature of fast reactors of 350 to 700°C, and the original properties of dispersion reinforced materials are not utilized. Therefore, the present invention aims to provide a dispersion-strengthened ferrite steel that has excellent properties required for a core member of a fast breeder reactor, that is, high-temperature strength and swelling resistance, and also has good ductility, toughness, and pipe formability. It has been done. <Means for Solving the Problem> That is, the first embodiment of the dispersion-strengthened ferrite steel for nuclear reactors having excellent ductility and toughness of the present invention has C: 0.1% or less, Si: 0.1% or less, Mn: 0.1
% or less, Cr: 12-20%, Mo + W = 0.1-4.4%,
O (excluding Y2O3 and TiO2 ): 0.01% or less, the balance consists of Fe and unavoidable impurities, and composite oxide particles of Y2O3 and TiO2 with an average particle size of 1000 Å or less are Y2O 3 + TiO2 : 0.1-1.0%, in molecular ratio
It is characterized by a ferrite single-phase or ferrite/martensite two-phase structure uniformly dispersed in the base (matrix) in the range of TiO 2 /Y 2 O 3 =0.5 to 2.0. Furthermore, a second embodiment of the dispersion-strengthened ferrite steel for nuclear reactors having excellent ductility and toughness according to the present invention has a weight %
C: 0.05% to 0.25%, Si: 0.1% or less, Mn: 0.1
% or less, Cr: 8 to 12% (however, 12% is not included),
Mo + W: 0.1 to 4.0%, O (excluding Y 2 O 3 and TiO 2 minutes): 0.01% or less, the remainder consisting of Fe and unavoidable impurities, and Y 2 O 3 with an average particle size of 1000 Å or less
Composite oxide particles made of TiO 2 are Y 2 O 3 + TiO 2 = 0.1
It is characterized by a tempered martensite single-phase structure uniformly dispersed in the matrix at a molecular ratio of TiO 2 /Y 2 O 3 of 0.5-1.0% and 0.5-2.0. The chemical composition of the dispersion-strengthened ferrite steel of the present invention and the reason for its limitation will be described below. this house
The most important point of the present invention is the combined addition of Y 2 O 3 and TiO 2 . Y 2 O 3 is the most important component that has the effect of improving creep rupture strength when uniformly dispersed in the matrix. However, Y 2 O 3 alone
It easily forms a complex oxide with a small amount of Si, Mn, etc. dissolved in the base, and becomes coarse. Also, Y 2
O 3 itself has poor compatibility with the matrix, and adding a large amount of Y 2 O 3 does not improve creep rupture strength, but rather reduces ductility and toughness. High strength can only be obtained by forming a stable composite additive of Y 2 O 3 and TiO 2 , Y 2 O 3 ·TiO 2 . The composite additive Y 2 O 3 ·TiO 2 , that is, Y 2 TiO 5 is formed by adding TiO 2 fine powder in the process of mechanically mixing the base alloy composition powder and Y 2 O 3 fine powder. Since Y 2 O 3 ·TiO 2 is more energetically stable than Y 2 O 3 alone, all of the Y 2 O 3 and TiO 2 react during mixing. It is also possible to use a Y 2 O 3 ·TiO 2 composite oxide prepared in advance. The above-mentioned prior patent also states that ``Ittria (Y 2 O 3 ) can be combined with other components in the composition, such as titanium, to form a phase such as Y 2 Ti 2 O 7. '' . However, with this method,
Since a composite oxide is formed by the reaction between Ti dissolved in the base and Y 2 O 3 particles, the composition of the composite oxide becomes non-uniform, with some having a high Ti concentration and others lacking Ti. Since such oxides are rather thermally unstable, if they are subjected to high-temperature softening treatment during pipe manufacturing, they will aggregate and become coarse, reducing the creep rupture strength. Further, excess Ti that does not react with Y 2 O 3 is precipitated as TiO 2 oxide alone. TiO 2 tends to coarsen during high-temperature use, causing a decrease in ductility. Moreover, once coarsening occurs, the ductility of the alloy cannot be restored even if high-temperature softening annealing is performed. In the present invention, the molecular ratio of Y 2 O 3 and TiO 2 is 0.5 to 2.0.
By reacting within this range, it is possible to uniformly disperse a complex oxide that is stable and has good compatibility with the base. In addition, since all the added TiO 2 and Y 2 O 3 react to form a stable composite oxide, high-temperature softening annealing restores the ductility to that before processing.
The amount of (Y 2 O 3 +TiO 2 ) is required to be at least 0.1% to improve high temperature strength. On the other hand, (Y 2 O 3 +
If the amount of TiO 2 ) added is increased, the creep rupture strength increases, but this effect is saturated at 1%, so the upper limit of the amount added is set at 1%. Furthermore, the oxide particle size is limited to 1000 Å or less since the effect of increasing creep strength is significantly reduced if the particle size exceeds 1000 Å. C is an important element that affects the stability of the structure, and its component range varies depending on the required characteristics. If ductility is important, the structure needs to be a single ferrite phase or a two-phase ferrite/martensite containing a small amount of tempered martensite.
In the first embodiment with a content of 12-20%, the amount of C is
It is necessary to keep it below 0.1%, preferably below 0.01%. In addition, by reducing the C content,
It is possible to increase the amount of solid solution of Mo and W in the matrix, and it is possible to increase the creep strength on the long-term side. On the other hand, if toughness is important, the structure needs to be a single phase of stable tempered martensite, and Cr
In the second embodiment with a content of 8 to 12%, the lower limit of the amount of C is 0.05%. This tempered martensitic structure is obtained by normalizing at 1000 to 1150°C and tempering at 700 to 800°C. As the C content increases, the amount of precipitated carbides (M 23 C 6 , M 6 C, etc.) increases and the high temperature strength increases, but if it is added in an amount greater than 0.25%, workability deteriorates. Therefore, the C content in this case is limited to 0.05 to 0.25%. The C content of Cr is 0.05% due to the balance with the C amount.
If less than 12% to 20%, ferrite phase, C
When the content is 0.05 to 0.25%, the tempered martensite phase can be stabilized in the range of 8 to 12%. Also, if it is less than 8%, the temperature will be high (600 to 700
decarburization resistance and corrosion resistance in sodium deteriorates. On the other hand, if it exceeds 20%, the toughness and ductility decrease, so the upper limit is set at 20%. Although Si is a necessary element as a deoxidizing agent, it easily reacts with Y 2 O 3 particles during use to form a composite oxide of Y 2 O 3 and SiO 2 . Since this composite oxide has a high coarsening rate, it reduces creep rupture strength. In addition, by reducing the Si content, the surface quality of the product can be improved and the amount of SiO 2 inclusions can be reduced, so it can be kept to 0.1% or less. Mn acts as a deoxidizing and desulfurizing agent and is an effective element for improving hot workability, but when added in large amounts, it forms a complex oxide with Y 2 O 3 that tends to become coarse like Si. % or less. Mo and W are important elements that form a solid solution in the alloy and improve high-temperature strength, and must be added in a total amount of 0.1% or more. If the amounts of Mo and W are increased, the creep rupture strength will improve due to solid solution strengthening, carbide precipitation strengthening (M 23 C 6 , M 6 C, etc.), and intermetallic compound precipitation strengthening, but Mo+W increases the strength by 4.0%. If it exceeds 4.0%, the amount of δ ferrite increases and the strength decreases, so the upper limit is set at 4.0%. In order to obtain particularly high strength, a combination of Mo and W with a Mo equivalent (Mo+1/2W) of 1.2 to 1.6% is preferable. O is an element that is inevitably contained in small amounts due to adsorption onto the raw material powder or oxidation, but if it exceeds 0.01%, the toughness will be significantly reduced. In addition, a small amount of Si and
Since inclusions are likely to form with Mn, the upper limit is set at 0.01%. <Example> The present invention will be described below with reference to Examples. Table 1 shows the chemical composition of the sample materials. In Table 1, steel types No. 1 to 3 are the invention steels () (corresponding to the first embodiment of the invention), and steel types No. 4 to 6 are the invention steels () (corresponding to the second embodiment of the invention). applicable),
Steel types No. 7 to 13 are comparative steels of the invention steel (), steel type 14
~20 is a comparative steel of the present invention steel (). The comparison steel contains important additives TiO 2 , Y 2 O 3 ,
(Mo+W), one in which the amount of C was out of the range of the present invention, or one in which Ti was added instead of TiO 2 was used. Among these, No. 13 corresponds to the alloy proposed in the prior patent. For each steel, element powder or alloy powder with an average particle size of 1 μm or less and oxide powder with an average particle size of 1000 Å or less are mixed to the desired composition, charged into a high-energy attritor, and then stirred in a high-purity argon gas atmosphere. Alloying was carried out mechanically. The rotation speed of the attritor is
200-300rpm, stirring time is 24-48hr. The obtained alloyed powder was vacuum sealed in a cylindrical SUS container without being exposed to air, and heated at 900 to 1200℃ for 8 to 10 minutes.
Hot extrusion was performed at an extrusion ratio of 15:1. After forging each hot pressed bar into a 10mm thick plate, 950
Normalized at ~1200℃, steel grades No. 1 to 3 and 7
~13 is as normalized, steel type No.4
~6 and 14~20 after normalizing, 750~
The test material was subjected to tempering heat treatment at 820℃. A plate-shaped tensile test piece of 2t x 6W x 30GL (mm) was taken from these test materials and subjected to a 650°C creep rupture test and a room temperature tensile test. Also, 5t×10W×55
mm (2 mm V notch) Charpy impact test pieces were taken and their impact properties were investigated. Furthermore, a 2 t x 6 W x 30 GL mm plate specimen was taken from a 10 (mm) thick plate material that had been 20% cold rolled and then annealed at 1200°C for 1 hour, and a room temperature tensile test was conducted to determine the change in tensile ductility after annealing. I looked into it. From those test results, the creep rupture stress at 650℃×10 3 hr, the tensile elongation at room temperature, the Charpy impact value at 20℃, and the tensile elongation after cold working + annealing are shown in Table 2.
are summarized in Both the invention steel () and () are compared to the comparative steel.
It can be seen that it has excellent creep rupture strength at 650°C, ductility at room temperature, Sharpy impact properties at 20°C, and ductility after cold working and annealing. Especially the invention steel ()
is excellent in ductility, and the steel of the present invention () is excellent in toughness. Among the comparison steels, steel types No. 7 and No. 14 have insufficient TiO 2 addition, so the dispersed oxide particles are unstable and the strength is low. Although steel grades No. 8 and No. 15 have relatively high strength,
Because the amount of TiO 2 added is excessive, the dispersed oxide particles aggregate and become coarse, resulting in low ductility and toughness. Steel types No. 9 and No. 16 are TiO2 - free, and steel types No. 10 and
In No. 17, no Y 2 O 3 was added, and the strength decreased significantly in both cases. This indicates that the dispersion is unstable when Y 2 O 3 and TiO 2 are added alone, and that the dispersion is stabilized and its strength increases only when Y 2 O 3 and TiO 2 coexist. Steel types No. 11 and 18 have Ti added instead of TiO 2 . Both have relatively high strength, but low ductility, especially after cold working and annealing. This indicates that in addition to reacting with Y 2 O 3 to form a composite oxide, Ti also precipitates as a single TiO 2 that tends to become coarse, so the TiO 2 particles become a stress concentration source and reduce ductility and toughness. There is. Steel types No. 12 and 19 have low strength due to insufficient (Mo+W) content, and although not shown in Table 2, the strength decrease is particularly large on the long-term side. In addition, since there are few solid solution elements, the base is not strengthened, and the toughness is lower than other steel types. Steel type No. 13 is an alloy proposed in a prior patent, but like steel types No. 11 and 18, Ti is added instead of TiO 2 , so it has lower ductility and toughness than the steel of the present invention. Steel type No. 20 has a lower C content than the steel of the present invention (). As a result, a large amount of ferrite phase appears, resulting in significantly low ductility and toughness.
【表】【table】
【表】【table】
【表】
＜発明の効果＞
以上に説明したように、本発明によれば長時間
優れた高温強度並びに延性及び靱性の良好な酸化
物分散強化型フエライト鋼が得られることから、
高速増殖炉炉心部材、特に燃料被覆管のような
650℃程度の高温で、しかも高い圧力下で使用さ
れる構造部材の長寿命化が達成できる。
また、延性と靱性が高く、軟化焼鈍による加工
材の延性回復が可能なことから、冷間加工を主体
とした製管法で経済性のある薄肉管の製管も可能
となる。[Table] <Effects of the Invention> As explained above, according to the present invention, an oxide dispersion strengthened ferrite steel with excellent long-term high temperature strength, good ductility and toughness can be obtained.
Fast breeder reactor core components, especially fuel cladding
It is possible to extend the lifespan of structural members that are used at high temperatures of around 650°C and under high pressure. In addition, since it has high ductility and toughness, and the ductility of the processed material can be recovered by softening annealing, it becomes possible to economically manufacture thin-walled pipes using a pipe manufacturing method that mainly involves cold working.
Claims (1)
Mn：0.1％以下、Cr：12〜20％、Mo＋Ｗ：0.1〜
4.0％、Ｏ（Y2O3およびTiO2分は除く）：0.01％以
下、残部Feおよび不可避不純物からなり、かつ
平均粒径1000Å以下のY2O3とTiO2による複合酸
化物粒子がY2O3＋TiO2＝0.1〜1.0％、分子比で
TiO2／Y2O3＝0.5〜2.0の範囲で基地に均一に分
散されているフエライト単相あるいはフエライ
ト／マルテンサイト２相組織であることを特徴と
する延性と靱性に優れた原子炉用分散強化フエラ
イト鋼。 ２ 重量％で、Ｃ：0.05〜0.25％、Si：0.1％以
下、Mn：0.1％以下、Cr：８〜12％（但し12％は
含まず）、Mo＋Ｗ：0.1〜4.0％、Ｏ（Y2O3および
TiO2分は除く）：0.01％以下、残部がFeおよび不
可避不純物からなり、かつ平均粒径1000Å以下の
Y2O3とTiO2による複合酸化物粒子がY2O3＋
TiO2＝0.1〜1.0％、分子比でTiO2／Y2O3＝0.5〜
2.0の範囲で基地に均一に分散されている焼戻し
マルテンサイト単相組織であることを特徴とする
延性と靱性に優れた原子炉用分散強化フエライト
鋼。[Claims] 1% by weight, C: 0.1% or less, Si: 0.1% or less,
Mn: 0.1% or less, Cr: 12~20%, Mo+W: 0.1~
4.0%, O (excluding Y 2 O 3 and TiO 2 ): 0.01% or less, the balance consists of Fe and unavoidable impurities, and the composite oxide particles of Y 2 O 3 and TiO 2 with an average particle size of 1000 Å or less are Y 2 O 3 + TiO 2 = 0.1-1.0%, in molecular ratio
Dispersion for nuclear reactors with excellent ductility and toughness characterized by a ferrite single phase or ferrite/martensite dual phase structure uniformly dispersed in the base with TiO 2 /Y 2 O 3 = 0.5 to 2.0. Reinforced ferrite steel. 2% by weight, C: 0.05 to 0.25%, Si: 0.1% or less, Mn: 0.1% or less, Cr: 8 to 12% (however, 12% is not included), Mo + W: 0.1 to 4.0%, O (Y 2 O3 and
(excluding TiO2 ): 0.01% or less, the remainder consisting of Fe and unavoidable impurities, and an average particle size of 1000Å or less
Composite oxide particles of Y 2 O 3 and TiO 2 are Y 2 O 3 +
TiO2 = 0.1 ~ 1.0%, molecular ratio TiO2 / Y2O3 = 0.5~
A dispersion-strengthened ferrite steel for nuclear reactors with excellent ductility and toughness, characterized by a tempered martensitic single-phase structure that is uniformly dispersed in the matrix in the range of 2.0.
Priority Applications (4)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP10229888A JPH01272746A (en) | 1988-04-25 | 1988-04-25 | Dispersion-strengthened ferritic steel for nuclear reactor excellent in toughness and ductility |
| US07/338,932 US4963200A (en) | 1988-04-25 | 1989-04-11 | Dispersion strengthened ferritic steel for high temperature structural use |
| GB8908952A GB2219004B (en) | 1988-04-25 | 1989-04-20 | Dispersion strengthened ferritic steel for high temperature structural use |
| FR8905316A FR2632659B1 (en) | 1988-04-25 | 1989-04-21 | FERRITIC STEEL REINFORCED BY DISPERSION FOR HIGH TEMPERATURE STRUCTURES |
Applications Claiming Priority (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP10229888A JPH01272746A (en) | 1988-04-25 | 1988-04-25 | Dispersion-strengthened ferritic steel for nuclear reactor excellent in toughness and ductility |
Publications (2)
| Publication Number | Publication Date |
|---|---|
| JPH01272746A JPH01272746A (en) | 1989-10-31 |
| JPH0518897B2 true JPH0518897B2 (en) | 1993-03-15 |
Family
ID=14323710
Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| JP10229888A Granted JPH01272746A (en) | 1988-04-25 | 1988-04-25 | Dispersion-strengthened ferritic steel for nuclear reactor excellent in toughness and ductility |
Country Status (1)
| Country | Link |
|---|---|
| JP (1) | JPH01272746A (en) |
Cited By (1)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| FR3007040A1 (en) | 2013-06-13 | 2014-12-19 | Japan Atomic Energy Agency Independent Administrative Corp | OXIDE DISPERSION REINFORCED TEMPERED STRENGTH MARTENSITIC STEEL WITH HIGH CORROSION RESISTANCE, TENACITY AND MECHANICAL PROPERTIES AT HIGH TEMPERATURE, AND PROCESS FOR PRODUCTION THEREOF |
Families Citing this family (5)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JPH05247587A (en) * | 1992-03-03 | 1993-09-24 | Nippon Steel Corp | High strength heat resistant steel plate reduced in radioactivation |
| FR2777020B1 (en) * | 1998-04-07 | 2000-05-05 | Commissariat Energie Atomique | PROCESS FOR MANUFACTURING A FERRITIC - MARTENSITIC ALLOY REINFORCED BY OXIDE DISPERSION |
| JP3753248B2 (en) * | 2003-09-01 | 2006-03-08 | 核燃料サイクル開発機構 | Method for producing martensitic oxide dispersion strengthened steel with residual α grains and excellent high temperature strength |
| JP2008546911A (en) * | 2006-06-20 | 2008-12-25 | フォルシュングスツェントルム カールスルーエ ゲゼルシャフト ミット ベシュレンクテル ハフツング | NUCLEAR FUEL MEMBER / FURTHER / MARTENSITE STEEL OR AUSTENITE STEEL COVER FOR NUCLEAR FUEL AND METHOD FOR POST-PROCESSING A FeCrAl Protective Layer Suitable On High Temperatures |
| KR20150104348A (en) * | 2014-03-05 | 2015-09-15 | 한국원자력연구원 | Ferrite/martensitic oxide dispersion strengthened steel with excellent creep resistance and manufacturing method thereof |
-
1988
- 1988-04-25 JP JP10229888A patent/JPH01272746A/en active Granted
Cited By (1)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| FR3007040A1 (en) | 2013-06-13 | 2014-12-19 | Japan Atomic Energy Agency Independent Administrative Corp | OXIDE DISPERSION REINFORCED TEMPERED STRENGTH MARTENSITIC STEEL WITH HIGH CORROSION RESISTANCE, TENACITY AND MECHANICAL PROPERTIES AT HIGH TEMPERATURE, AND PROCESS FOR PRODUCTION THEREOF |
Also Published As
| Publication number | Publication date |
|---|---|
| JPH01272746A (en) | 1989-10-31 |
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