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castability of gray cast iron in lost foam process

2010-09-15T06:44:00.000-07:00

عوامل موثر در سیالیت چدن خاکستری در ریخته گری لاست فوم

ميريس سيد شريعت دوست ، واحد نجاتی مازگر

شرکت ریخته گری تراکتورسازی ایران-تبریز

چکیده:

در رابطه با تاثیر برخی عوامل متالورژیکی از قبیل دمای ذوبریزی ، ترکیب شیمیائی و ... بر روی سیالیت مذاب در آلیاژهای مختلف ریخته شده در ماسه قبلا مقالاتی منتشر شده است.این در حالی است که در رابطه با این موضوع در ریخته گری لاست فوم اطلاعات چندانی در اختیار نمی باشد.لذا در این مقاله بر آن شدیم تا در کنار بررسی تاثیر متغیرهائی چون دمای بارریزی،درصد کربن معادل و ضخامت مقطع ریختگی تاثیر عواملی چون دانسیته و ضخامت رنگ مدل فومی را نیز روی سیالیت مذاب چدن خاکستری در روش لاست فوم مورد بحث و بررسی قرار دهیم.

در تحقیق حاضر به روش لاست فوم نمونه های اسپیرال با ضخامت مقطع ،دانسیته و ضخامت پوشان مختلف از ذوب خاکستری با درصد کربن معادل 4.3- 3.8 % در دامنه دمائی C° 1430-1380 ریخته شدند.نتایج حاصل از آزمایشات بیانگر افزایش سیالیت مذاب با افزایش درصد کربن معادل ، دمای ذوبریزی و ضخامت پوشان می بود و نیز با افزایش دانسیته مدل فومی سیالیت مذاب کاهش نشان می داد. نتیجه دیگر وجود رابطه خطی بین سیالیت مذاب و ضخامت مقطع می بود .

کلید واژه: سیالیت،نمونه اسپیرال،قابلیت ریخته گری،ضخامت مقطع

مقدمه:

جریان مذاب در طول ذوبریزی و پر شدن قالب پایه و اساس علم ریخته گری به شمار می رود.قابلیت مذاب در پر کردن قالب قبل از بروز انجماد را سیالیت مذاب می نامند که از خواص ذاتی و حیاتی مذاب در ریخته گری به شمار می رود.طبق تعریف سیالیت به طول مسافتی از یک قالب کانالی با ضخامت کم اطلاق می گردد که فلز مذاب قبل از بروز انجماد قادر به طی آن می باشد و بر حسب سانتی متر یا اینچ بیان می گردد.امروزه با توجه به افزایش پیچیدگی در قطعات ریختگی و وجود مقاطع با ضخامت بسیار کم بخصوص در قطعات خودرو ،درک صحیح قابلیت ریخته گری و حصول سیالیت لازم در مذاب امری اجتناب ناپذیر می باشد.این امر در ریخته گری به روش لاست فوم اهمیت دوچندان پیدا می کند که در این روش بالذاته سیالیت مذاب و و قابلیت ریخته گری بخاطر حضور مدل فومی و توپر بودن قالب کاهش می یابد(شكل 1).

همانطور كه به وضوح در شكل 1 ديده مي شود وجود مذدل فومي و محصولات مايع و گازي حاصل از سوختن آن در طول ذوبريزي باعث افت حرارت مذاب و كاهش سياليت مي گردد.لازم به ذكر است مقدار اين محصولات حاصل از سوختن فوم بيشتر بستگي به دماي مذاب و دانسيته مدل فومي دارد.براي مثال در ريخته گري آلومنيوم با دماي 750 درجه سانتي گراد مقدار محصولات گازي 40 سانتي متر مكعب به ازاي هر گرم فوم مي باشد كه رقم قابل ملاحظه اي مي باشد و مي تواند مانع پيشروي سريع جبهه مذاب در درون قالب گردد.

سیالیت اگرچه از خواص ذاتی مذاب می باشد و بيشتر تحت تاثير ويسكوزيته مذاب ،گرماي نهان ذوب،كشش سطحي مذاب ،دامنه انجماد و دانسيته فلز مي باشد ولی عوامل متعدد ديگري نيز روی آن تاثیر گذار می باشند .در ریخته گری با قالبهای ماسه ای ترکیب شیمیائی آلیاژ و در رابطه با چدنها درصد کربن معادل تاثیر به سزائی روی سیالیت مذاب دارد.عامل مهم دیگر که اثر آن حتی بیشتر از درصد کربن معادل است،دمای ذوبریزی می باشد.سیالیت مذاب با گرادیان حرارتی و سرعت انتقال حرارت از مذاب ارتباط تنگاتنگی دارد و بدیهی است با افزایش سرعت سرد شدن و انجماد از سیالیت آلیاژ کاسته می شود.به همین دلیل عامل موثر دیگر در سیالیت مذاب ضخامت مقطع ریختگی می باشد.در ریخته گری لاست فوم علاوه بر موارد فوق عوامل دیگری نیز روی سیالیت مذاب اثر گذار می باشند.از جمله مهم ترین این عوامل می توان به ضخامت پوشان اعمال شده روی مدل فومی ،جنس پوشان و میزان انتقال حرارتی در آن،دانسیته مدل فومی و نحوه پر شدن قالب می باشد.

با توجه به موارد اشاره شده به وضوح دیده می شود که سیالیت مذاب مستقیما بر روی کیفیت قطعه ریختگی اثرگذار می باشد و در صورت عدم وجود سیالیت لازم در مذاب احتمال بروز عیوبی از قبیل کم آمد، اتصال سرد، حفرات گازی،نشتی و ... در قطعات ریختگی بالا خواهد بود.هدف اصلی از انجام این تحقیق تبیین رابطه سیالیت با عوامل فوق در ریخته گری چدن خاکستری به روش لاست فوم می باشد که به شرح زیر انجام پذیرفت.

روش انجام آزمایشات:

همانگونه که اشاره شد هدف از انجام این تحقیق بررسی عوامل موثر بر سیالیت چدن خاکستری در روش لاست فوم می باشد. عوامل مورد نظر و دامنه در نظر گرفته شده برای این عوامل که پایه و اساس آزمایشات در تحقیق حاضر را تشکیل می دهند ،به شرح زیر در جدول 1 آورده شده است.

جدول 1 ) پارامترهای موثر و دامنه مورد نظر آنها در آزمایشات

دمای ذوبریزی درصد کربن معادل ضخامت مقطع ریختگی ضخامت پوشان مدل فومی دانسیته مدل فومی

C° 1430- 1380 4.3-3.8% mm2،1 و 3 mμ 900-300 gr/lit 35-25

در ابتدا برای شروع آزمایشات مدل فومی اسپیرال(حلزونی شکل جهت تست سیالیت) طراحی و ساخته شد . در این رابطه شایان ذکر است با توجه به گستردگی موضوع در فرایند تولید مدل های فومی از بیان کامل جزئیات غیر مرتبط با موضوع مقاله اجتناب گردید و اطلاعات مورد نیاز در این زمینه به شرح زیر می بود:

الف) ذرات فوم اولیه خام :

• جنس: پلی استایرن منبسط شونده (EPS)

• دانه بندی فوم اولیه خام : 0.03-0.05 mm

• دانسیته اولیه: 0.53-0.65 g/lit

• میزان مواد فرار محتوی: ≥ 7.5%

ب) مرحله پیش انبساطی :

برای این منظور از یک دستگاه Pre-expander آزمایشگاهی استفاده به عمل آمد.ابتدا به میزان 3.1 پوند از فوم خام (فوق الذکر) در دستگاه ریخته و با دمش بخار با شرایط زیر انبساط اولیه صورت گرفت:

• زمان بخار دهی : 10-50 ثانیه (بسته به دانسیته نهائی مورد نیاز در مدلهای آزمایشی برای تست سیالیت )

• فشار بخار : 2 bar

• دانسیته فوم منبسط شده : gr/lit 30-20

ج) مرحله پیرسازی:

در این مرحله فوم های حاصل از مرحله فوق خشک شده و به مدت 12-8 ساعت بلا استفاده باقی ماندند تا به پایداری لازم برسند.در این مرحله رطوبت و پنتان باقی مانده از ذرات فوم خارج می شود.

د)ساخت مدل:

در این مرحله فوم های پیرسازی شده به شرح زیر جهت تولید مدل های اولیه تست سیالیت مورد استفاده قرار گرفتند:

• بستن لنگه های قالب آلومنیومی از پیش ساخته شده با ونت های دمشی کافی در سطوح آن

• شستن قالب با بخار جهت تست باز بودن ونت ها

• بررسی اجکتورها ( وکیومی برای در آوردن مدل های تولیدی از قالب)

• شوت کردن وزن معین فوم منبسط شده و پیر سازی شده به قالب

• تزریق بخار با فشار و دمای معین به درون قالب

• تزریق آب سرد برای خودگیری مدل فومی و سهولت در خارج سازی آن

• باز کردن لنگه های قالب و درآوردن مدل با سيستم مكشي

ه) مونتاز و خوشه سازی :

حال پس آماده شدن تکه های مدل و راهگاهها برای اتصال و چسباندن آنها از نوعی چسب گرم ساخت کشور آلمان با مشخصات زیر استفاده به عمل آمد:

و) پوشان دهی:

در این مرحله خوشه های تولید شده به روش غوطه وری و با استفاده از پوشان با مشخصات زیر رنگ آمیزی گردید.لازم به ذکر است با توجه به اینکه ضخامت رنگ روی مدل از پارامترهای آزمایشی در تحقیق حاضر به شمار می رود لذا برای اعمال ضخامت مورد نظر از پوشان دهی چند مرحله ای استفادده به عمل آمد و در هر مرحله ضخامت تقریبا 0.3 میلیمتر روی مدل فومی ایجاد می گردید.

• دانسیته رنگ در دمای محیط: 1.75 gr/cm3

• اسیدیته (PH) : 7.5

• ویسکوزیته : 13-14Sec Cup ,DIN 6mm

• ضخامت رنگ روی مدل: 0.3 میلیمتر در هر بار فروبری و 0.9-0.7 میلیمتر با دو بار فروبری در پوشان.

ز) ذوبریزی:

برای انجام آزمایشات و ذوبریزی ،خوشه های فومی ساخته شده براي تست سياليت (شكل 2) در داخل درجه ها ی حاوی ماسه خشک سیلیسی با عدد ریزی AFS 33-37 (پیوست 2) قرار داده شد و از ماسه پر گردید .در ادامه با اعمال ویبره به ميزان 1.2g و به مدت زمان 45-30 ثانيه و پر كردن قالبها از ماسه در طول ويبراسيون ،قالبها آماده برای ذوبریزی گردیدند.

(a)

b) ) (c)

شکل 2) تصاویری از(a) مدل اسپیرال فومی ،b) ) اتصال و مونتاژ حوضچه و ملحقات آن به مدل اسپیرال فومی و (c) نمونه هائی از مدل های اولیه حلزونی با دانسیته های متفاوت(قبل از برش و مونتاژ)

مذاب مورد استفاده در این تحقیق از جنس خاکستری می بود که با درصد های متفاوت کربن معادل و در دماهای مختلف ذوبریزی (مطابق جدول 2 ) در قالبها ریخته شد. (جزئیات مربوط به انالیز مذاب در نمونه های ازمایشی در پیوست 1 اورده شده است.

جدول 2) شرایط انجام آزمایشات و کد گذاری نمونه ها

دمای ذوبریزی

C°5±

کربن معادل

% ±0.05

نمونه های سری1 ضخامت پوشان

μ10± دانسیته مدل

gr/lit2±

نمونه های سری2 ضخامت

مقطع

mm

نمونه های سری3

1380 3.8 A

300 25 J

1

S

4.05 B 30 K

4.3 C 35 L

1405 3.8 D

600 25 M

2

T

4.05 R 30 N

4.3 F 35 O

1430 3.8 G

900 25 P

3

U

4.05 H 30 Q

4.3 I 35 R

با توجه به تعدد پارامترها در این تحقیق که در جدول 2 نیز نشان داده شده است جهت هدفمند کردن جمع بندی نتایج ،نمونه های آزمایشی را به سه سری به شرح زیر تقسیم بندی کردیم:

نمونه های سری 1)

همانطور که در جدول 2 نشان داده شده است نمونه های سری 1 شامل نمونه های آزمایشی مربوط به تغییرات سیالیت با درصد کربن معادل در دماهای ذوبریزی مختلف می باشد . در این تحقیق درصد کربن معادل 3.8،4.05 و4.3 در نظر گرفته شده است و برای هر کدام از این درصد ها نمونه هائی در سه دمای بارریزی 1380،1405 و 1430 درجه سانتی گراد ریخته شدند و لذا این سری از نمونه ها شامل 9 عدد نمونه خواهد بود که از A تا I نشان داده شده است.

نمونه های سری 2)

همانطور که در جدول 2 نشان داده شده است نمونه های سری 2 شامل نمونه های آزمایشی مربوط به تغییرات سیالیت با دانسیته مدل فومی در ضخامت های مختلف پوشان مدل می باشد . در این تحقیق دانسیته مدل فومی 25-35gr/lit در نظر گرفته شده است و برای هر کدام از این شرایط نمونه هائی در ضخامت های مختلف پوشان به میزان mμ 300-900 تهیه و آزمایش شدند و لذا این سری از نمونه ها نیزشامل 9 عدد نمونه خواهد بود که از J تا R نشان داده شده است.

نمونه های سری 3)

در این سری از نمونه ها همانطور که در جدول 2 نیز نشان داده شده است تغییرات ضخامت مقطع قطعه با سیالیت مذاب مورد بحث و بررسی قرار گرفت است.در این حالت نمونه های با ضخامت مقطع 1،2و 3 میلیمتر آزمایش شدند و سیالیت مذاب در هر کدام از این شرایط تست و محاسبه گردید. این سری از نمونه ها نیزشامل 3 عدد نمونه می بود که از S تا U نشان داده شده است.

بحث و نتیجه گیری:

در جدول 3 نتایج حاصل از تست سیالیت به عمل آمده بر روی نمونه های سری 1 با ضخامت مقطع 2 میلی متر نشان داده شده است.لازم به ذکر است در این سری از نمونه ها ضخامت مقطع نمونه ها 2 میلی متر، دانسیته مدل های فومی : g/lit 1 ± 30 و ضخامت پوشان روی مدل: mμ 50 ± 600 و ثابت می بود.

جدول 3) نتایج تست سیالیت مربوط به نمونه های سری 1

سیالیت مذاب (Cm) درصد کربن معادل

% ±0.05 دمای ذوبریزی

°C ±5 کد نمونه

41 3.8 C°1380 A

47 4.05 C°1380 B

52 4.3 C°1380 C

43 3.8 C°1405 D

52 4.05 C°1405 E

59 4.3 C°1405 F

56 3.8 C°1430 G

63 4.05 C°1430 H

76 4.3 C°1430 I

همانطور که در جدول فوق ملاحظه می گردد با افزایش دما و نیز درصد کربن معادل ،سیالیت مذاب افزایش می یابد.حال در این قسمت جهت تحلیل بهتر موضوع در شکل 3 نمودار تغییرات سیالیت با درصد کربن معادل در دماهای مختلف بارریزی نشان داده شده است.همانطور که ملاحظه می گردد شدت افزایش سیالیت با افزایش دما نسبت به افزایش درصد کربن معادل بیشتر می باشد.به عبارتی سیالیت مذاب با درصد کربن معادل پائین در دمای بالاتر بیشتر از سیالیت مذاب با کربن معادل بالا در دمای بارریزی پائین می باشد.لذا نتیجه می گیریم برای ریخته گری قطعات با ضخامت مقطع کم از قبیل بلوک سیلندر ها لازم است ضمن بالا گرفتن نسبی درصد کربن معادل ضروری است دمای بارریزی تا حد ممکن بالاتر انتخاب گردد.

شکل 3) تغییرات سیالیت مذاب با درصد کربن معادل در دماهای مختلف ذوبریزی

در ادامه آزمایشات تاثیر ضخامت پوشان و دانسیته مدل فومی روی سیالیت مذاب مورد بررسی قرار گرفت.برای این منظور مدل های اسپیرال با دانسیته 30،25 و 35 گرو بر لیتر تولید گردید و روی هر کدام از این مدل ها پوشاندهی با ضخامت های مختلف 600،300 و 900 میکرون صورت پذیرفت .بدین ترتیب 9 عدد خوشه فومی بدست آمد که مطابق جدول 2 با حروف J تا R نشان داده شده است.لازم به ذکر است در نمونه های مربوط به این سری از آزمایشات درصد کربن معادل 4±0.05% و دمای ذوبریزی 1400±5°C و در ضمن ضخامت نمونه ها نیز 2 میلی متر بود.

جدول 4) نتایج تست سیالیت مربوط به نمونه های سری 2

سیالیت مذاب (Cm) دانسیته مدل

gr/lit ±2 ضخامت پوشان

m μ ±10 کد نمونه

41 25 300 J

52 30 300 K

60 35 300 L

44 25 600 M

54 30 600 N

63 35 600 O

45 25 900 P

58 30 900 Q

65 35 900 R

همانطور که در جدول فوق ملاحظه می گردد با افزایش ضخامت پوشان و نیز کاهش دانسیته آن سیالیت مذاب افزایش می یابد ولی جهت تحلیل بهتر موضوع در شکل 4 نمودار تغییرا ت سیالیت با این پارامترها را بررسی خواهیم کرد.

شکل 4) نمودار تغییرات سیالیت مذاب با ضخامت پوشان و دانسیته مدل فومی

همانطور که در شکل فوق به وضوح دیده می شود و قبلا نیز اشاره شد با افزایش ضخامت پوشان و نیز کاهش دانسیته مدل فومی سیالیت مذاب افزایش می یابد و در این میان شدت و نرخ تغییرات سیالیت با دانسیته مدل بیشتر از ضخامت پوشان می باشد و کاهش دانسیته باعث افزایش فاحش در سیالیت مذاب می شود و در ضمن میزان تغییرات سیالیت در دانسیته های پائین کمتر از دانسیته های بالاتر می باشد. لذا برای ریخته گری قطعات پیچیده با ضخامت کم ضروری است که دانسیته مدل فومی تا حدی پائین در نظر گرفته شود.

در مرحله بعدی آزمایشات تاثیر ضخامت مقطع روی سیالیت (نمونه های سری سوم ) مورد بحث و بررسی قرار گرفت. در این مرحله از آزمایشات ضخامت مقطع نمونه 1،2 و3 میلی متر وسایر شرایط به شرح زیر ثابت در نظر گرفته شد:

دمای ذوبریزی :5± 1405 درجه سانتی گراد ، درصد کربن معادل مذاب : 0.05%±4 ، دانسیته مدل های فومی : g/lit 1 ± 30 و ضخامت پوشان روی مدل: mμ 50 ± 600 .

در جدول 5 نتایج مربوط به آزمایشات این سری نمونه ها آورده شده است.

جدول 5) نتایج تست سیالیت مربوط به نمونه های سری 3

سیالیت مذاب (Cm) ضخامت مقطع نمونه (mm) کد نمونه

41 1 S

54 2 T

84 3 U

همانطور که قبلا اشاره شد افزایش ضخامت مقطع در نمونه ریختگی، سیالیت مذاب نیز افزایش می یابد.نتایج آزمایشات بیشتر در این زمینه بیانگر وجود ارتباط خطی بین سیالیت مذاب با ضخامت مقطع می باشد. حال با توجه به افت سیالیت در مقاطع با ضخامت کم برای جلوگیری از بروز عیوب مرتبط ضروری است کربن معادل ،دمای بارریزی و ضخامت پوشان بالاتر و دانسیته مدل فومی کمتر انتخاب گردد.

نتایج:

با توجه به بحث و بررسی های صورت پذیرفته نتایج حاصل از آزمایشات به عمل آمده در این تحقیق را می توان به شرح زیر خلاصه نموند:

1. سیالیت و قابلیت ریخته گری مذاب از پارامترها و ویژگی های مهم در فرایند ریخته گری به شمار می رود که روی کیفیت قطعه ریختگی نقش حیاتی را ایفا می نماید و لذا مقدار آن بایستی از حد معینی کمتر نباشد.

2. از جمله مهم ترین عوامل موثر در سیالیت قطعات ریخته شده به روش لاست فوم می توان به دمای ذوبریزی، ضخامت مقطع،دانسیته مدل فومی ، درصد کربن معادل و ضخامت پوشان به ترتیب اهمیت و تاثیرشان اشاره کرد .

3. با افزایش دمای ذوبریزی ،ضخامت مقطع و رنگ مدل فومی و درصد کربن معادل سیالیت مذاب افزایش می یابد و می توان قطعات با ضخامت نازک و پیچیده را ریخته گری نمود.در مقابل با افزایش دانسیته مدل فومی سیالیت مذاب کم می شود.در این میان بین ضخامت مقطع و سیالیت مذب یک رابطه خطی برقرار می باشد

4. در هر سیستم آلیاژی آلیاژ با ترکیب یوتکتیک بدلیل داشتن حداکثر مقدار سیالیت مناسب ترین آلیاژ برای ریخته گری می باشد.حال چنانچه اطلاع داریم در چدن ها ترکیب یوتکتیک در کربن معادل تقریبا 4.3 درصد قرار می گیرد ولذا در آلیاژهای با درصد کربن معادل بالاتر از 4.3 درصد (اکثرا چدن داکتیل) با افزایش کربن معادل و دور شدن از ترکیب یوتکتیک ،سیالیت مذاب کاهش می یابد ولی در چدنهای خاکستری که عموما درصد کربن معادل کمتر از 4.3 درصد می باشد با افزایش درصد کربن معادل و نزدیک شدن به ترکیب یوتکتیک سیالیت آلیاژ افزایش می یابد.

پیوست 1)

آنالیز نمونه های ریخته شده در این تحقیق که نتایج نهائی آن (کربن معادل) در جدول 2 آورده شده بود به شرح می بود:

جدول a) ترکیب شیمیائی نمونه های آزمایشی

Al Mg Sn Cu S Mn Si C CE کد نمونه ها

0.001 0.010 0.078 0.66 0.111 0.32 1.702 3.266 3.839 A

0.002 0.000 0.088 0.67 0.092 0.39 1.640 3.463 4.016 B

0.002 0.001 0.087 0.60 0.082 0.34 1.565 3.751 4.279 C

0.001 0.010 0.080 0.74 0.109 0.28 1.650 3.258 3.813 D

0.001 0.008 0.085 0.65 0.081 0.35 1.422 3.535 4.015 E

0.002 0.006 0.073 0.76 0.109 0.29 1.574 3.766 4.296 F

0.002 0.009 0.086 0.65 0.090 0.39 1.525 3.321 3.835 G

0.001 0.013 0.087 0.76 0.088 0.28 1.531 3.497 4.013 H

0.002 0.016 0.085 0.66 0.091 0.27 1.619 3.753 4.298 I

0.001 0.013 0.087 0.55 0.071 0.35 1.588 3.424 3.998 J

0.003 0.000 0.086 0.57 0.062 0.38 1.588 3.424 3.959 K

0.002 0.001 0.087 0.60 0.072 0.35 1.575 3.400 3.931 L

0.000 0.019 0.088 0.64 0.079 0.35 1.549 3.397 3.949 M

0.000 0.000 0.086 0.54 0.081 0.43 1.67 3.365 3.957 N

0.001 0.001 0.084 0.56 0.069 0.40 1.710 3.523 4.099 O

0.001 0.001 0.083 0.65 0.080 0.34 1.696 3.382 3.953 P

0.002 0.010 0.086 0.47 0.078 0.30 1.546 3.450 3.970 Q

0.001 0.001 0.086 0.56 0.089 0.33 1.700 3.466 4.038 R

0.001 0.001 0.085 0.64 0.090 0.32 1.581 3.409 3.952 S

0.002 0.010 0.090 0.58 0.098 0.41 1.735 3.457 4.041 T

پیوست 2)

در این تحقیق نتایج حاصل از نحوه توزيع و دانه بندي ماسه قالبگيري مورد استفاده به شرح زیر می بود.

منابع و مراجع:

1- THE EFFECT OF MOLD MATERIALS ON SOLIDIFICATION,

MICROSTRUCTURE AND FLUIDITY OF A356 ALLOY IN LOST

FOAM CASTING by, Ramin Ajdar

2- I O. R.W.Monroe, band able Pattern Casting, AFS, Des Plaines, Illinois, U.S.A,

1992, pp. 96-97.

3- E L Kotzin, Metal Castinci & Moldina Processes, AFS, Des Plaines, lllinois,

USA, 1981, pp.149-153.

4- S.Shivkumar, "Casting Characteristics of Aluminum Alloy in the EPC Processn

AFS Transactions, Vol.101,1993, pp.513-5 t 8.

5- A.T.Spada, "Core Competency lncludes Lost Foam", Modem Casting, Vol.

91, May 2001, pp. 27-30.

6- J .Campbell , Castinas, Butferworth-Heinemann LM, Oxford, England, 1 991, p.

143.

تشكر و قدرداني:

در اين اين قسمت دارد از زحمات بيدريغ جناب آقايان مهنسين قديمي و بندرچي و سركار خانم اميرخاني كه بدون حمايتهاي آنها امكان ارائه اين تحقيق و كارهاي مشابه مقدور نمي شد كمال تشكر و امتنان را دارم .در نهايت نيز از راهنمائيهاي ارزنده جناب آقايان دكتر دوامي و ورهرام كه همواره در زنگي چراغ راه من بوده است كمال امتنان را داشته و آرزوي سلامتي روزافزون برايشان را از درگاه ايزد منان خواهانم.

production of ascast ferritic ductile iron

2010-09-15T06:42:00.001-07:00

پارامترهاي كنترلي در توليد چدن داكتيل فريتي در حالت سياه تاب

مهندس واحد نجاتي مازگر

كارشناس ارشد متالورژي- شركت ريخته گري تراكتورسازي ايران- تبريز

چكيده:

در اين تحقيق اثر پارامترهائي از قبيل مواد شارژي،شرايط ذوب ،ضخامت مقطع،روش جوانه زني و كروي كردن روي تركيب شيميائي مذاب و ريزساختار زمينه در قطعات چدني داكتيل مورد بحث و بررسي قرار گرفته است.نتايج نشان مي دهد كه افزايش دما و زمان نگهداري مذاب در كوره باعث كاهش درصد كربن و افزايش مقدار سيليسيم محتوي مذاب كوره مي گردد.از طرف ديگر سرعت و شدت عمليات تلقيح منيزيم و كروي كردن گرافيتها روي مقدار كربن و منيزيم باقي مانده در مذاب تاثير گذار مي باشد.در اين تحقيق مشاهده گرديد كه براي توليد چدن داكتيل فريتي در حالت سياه تاب نياز به استفاده از مقادير زيادي شمش در تركيب مواد شارژي كوره مي باشد و در صورت كم بودن ضخامت مقطع بايستي از عمليات جوانه زني چند مرحله اي نيز استفاده به عمل آيد.

مقدمه:

در طول سالهاي اخير چدن داكتيل به به خاطر داشتن قابليت ريخته گري بالا ،هزينه توليد پائين و خواص مكانيكي عالي يك ماده مهم مهندسي به شمار مي رود.اين ماده به دليل دارا بودن دامنه گسترده اي از خواص در كاربردهاي زيادي مورد استفادده قرار مي گيرد و انواع عمليات حرارتي شده آن اخيرا جايگزين مناسبي براي فولاد فورج شده است.خواص حاصله در چدن داكتيل عمدتا به ريزساختار آن يعني زمينه و شكل،اندازه و نحوه توزيع گرافيتهاي كروي در آن بستگي دارد.از عوامل موثر و تعيين كننده در نوع زمينه و شكل و اندازه گرافيتها نيز مي توان به تركيب شيميائي،عمليات گوگرد زدائي،تلقيح منيزيم،جوانه زني و فاصله زماني بين اين عملياتها اشاره كرد.

توليد چدن داكتيل و خواص حاصل از آن به پارامترهاي متعددي مانند متالورژيكي،تكنولوژيكي،هدايت حرارتي و طراحي بستگي دارد.اولين گام در توليد چدن داكتيل انتخاب صحيح مواد شارژي مي باشد. منگنز و كروم در بيشترين تاثير را روي خواص مكانيكي دارند و منبع ورود آنها به تركيب شيميائي مذاب ،قراضه هاي فولادي و برگشتي هاي جدني مي باشند.توصيه مي شود مقدار اين دو عنصر زير 0.1 درصد در مواد شارژي بوده باشند.ولي متاسفاده در رابطه با منگنز اكثر قراضه هاي فولادي بيش از 0.5 درصد منگنز دارند وكنترل آن تا حدودي مشكل مي باشد.از طرف ديگر مقدار قراضه هاي فولادي در تركيب مواد شارژي بايستي بتا حدي باشد كه از بروز كاربيد در قطعه ريختگي جلوگيري شود و اين موضوع بخصوص در چدنهاي داكتيل با زمينه فريتي از اهميت بالاتري برخوردار مي باشد.

علاوه بر موارد ذكر شده مواد شارژي روي اندازه متوسط كره هاي گرافيتي نيز مستقيما اثرگذار مي باشد.براي مثال اگر مقدار قراضه فولادي در تركيب مواد شارژي از 50 درصد بيشتر باشد قطر متوسط كره هاي گرافيتي حدود 34μخواهد بود و اگر مقدار آن به 30 درصد كاهش يابد قطر متوسط كره هاي گرافيتي نيز تا μ57 افزايش خواهد يافت.از طرف ديگر با افزايش مقدار قراضه فولادي در مواد شارژي زمينه ريزساختاري نيز به سمت پرليتي سوق مي يابد.

از ديگر عواملي كه ساختار گرافيت را شديدا تحت تاثير قرار مي دهد درصد كربن محتوي مذاب مي باشد .اگر مذاب پايه از كربن كافي برخوردار نباشد در آن صورت گرافيتها به شكل فشرده ظاهر خواهند شد.درصد كربن همچنين روي زمينه نيز موثر مي باشد به گونه اي كه با افزايش درصد كربن معادل و نيز با كاهش نسبت كربن به سيليسيم ميزان فريت در زمينه افزايش مي يابد و زمينه به سمت پرليتي شدن ميل مي كند.مقدار عناصر آلياژي در تركيب شيميائي مواد شارژي به شدت نوع زمينه را تحت تاثير قرار مي دهد به گونه اي كه براي حصول زمينه فريتي لازم است شرايط زير در تركيب مواد شارژي برقرار باشد: Sn≤0.001%, As≤0.02%, (V+Mo)≤0.1%, Cr≤0.04

همانطور كه مي دانيم منيزيم متداول ترين عنصر براي كروي كردن گرافيتها در چدن مي باشد و در اين ميان نحوه افزودن منيزيم(روش داكتيل سازي ) و نوع آلياژ مصرفي از اهميت بالائي برخوردار مي باشد و انتخاب آنها به شرايط و امكانات واحد ريخته گري بستگي دارد.

فرايند ديگري كه در پروسه توليد چدن داكتيل از اهميت به سزائي برخوردار مي باشد عمليات جوانه زني مي باشد كه مي تواند به طرق مختلفي از قبيل افزودن ماده جوانه زا به جريان مذاب،به پاتيل و يا به قالب صورت پذيرد و بسته به خواص مورد نظر و امكانات كارگاهي مي توان از اين عمليات به صورت تكي يا تركيبي استفاده به عمل آورد.متداول ترين ماده براي جوانه زني چدنها آلياژ فروسيليسيم همراه با مقادير كمي از عناصر آلياژي ديگر از قبيل باريم،استرانسيم،زيركونيوم،كلسيم،آلومنيوم و ... مي باشد كه هر كدام از اين عناصر براي نيل به اهداف و خواص خاصي مورد استفاده قرار مي گيرند.

حال همانطور كه به وضوح ديده مي شود و اجمالا نيز اشاره شد پارامترهاي متعددي روي خواص نهائي چدن هاي داكتيل تاثير گذار مي باشند و در اين تحقيق روي متغيرهاي موثر در توليد چدنهاي داكتيل با زمينه فريتي در حالت سياه تاب صحبت خواهد شد.

آزمايشات عملي:

براي تهيه مذاب در اين تحقيق از يك دستگاه كوره القائي 14 تني ساخت شركت ABB كه با فركانس شبكه كار مي كرد استفاده به عمل آمد.تركيب مواد شارژي مورد استفاده در تهيه مذاب در جدول 1 آورده شده است.همچنين در جدول 2 تركيب شيميائي مذاب پايه داكتيل نشان داده شده است.براي گوگردزدائي مذاب از كاربيد كلسيم به مقدار لازم در داخل كوره استفاده به عمل آمد.پس از آمده شدن مذاب كم گوگرد براي تهيه مذب داكتيل جهت انجام آزمايشات از روش تانديش كاور استفاده به عمل آمد .در شكل 1 تصويري از پاتيل تانديش دار مورد استفاده در اين تحقيق براي داكتيل سازي مذاب نشان داده شده است.مواد مورد استفاده براي داكتيل سازي آلياژ FeSiMg با 6 درصد منيزيم ساخت شركت LAMET مي بود.

Controlling cast iron gas defects

2010-05-15T02:10:00.000-07:00

Controlling cast iron gas defects.

Gas defects stem from internal chemical reactions, external mechanical pressures or gas solubility limits of the molten metal.

Gas defects can be classified by their gas source, but a literature search reveals the terminology for gas defects is often inconsistent. The term "blowhole" is usually associated with a larger cavity than that described as a "pinhole."

The distinction created by the following terminology seeks to differentiate between the various gas defects.

Reaction

Gas Holes: caused by the interaction between carbon (C) dissolved in iron and a nonmetallic inclusion to generate carbon monoxide (CO). Described as an endogenous (internally generated) gas defect, it is usually a blowhole from a reaction with slag carried into the mold. These reaction voids may be located throughout a section of a casting but are normally near the cope surface. They may be large, small or irregular with either a smooth or rough shape.

Mechanical Gas Holes: caused when gas is either entrained during pouring or is forced into the shape from the mold or core after pouring as depicted in the schematic in Fig. 2. They are referred to as exogenous (externally caused) gas holes.

The required pressure is defined by liquid metal head and surface tension and bubble radius. A gas hole forms when the gas pressure, at any point, exceeds the head pressure and is influenced by the mold's ability to release the gases generated in the mold (permeability).

The defect is generally large, but may have a small entry hole. It can be located at the source of the trapped gas but may rise toward the thermal center of the casting, where the occurrence of shrinkage cavities can make differentiation between the two defects difficult

Gas defects stem from internal chemical reactions, external mechanical pressures or gas solubility limits of the molten metal.

Gas defects can be classified by their gas source, but a literature search reveals the terminology for gas defects is often inconsistent. The term "blowhole" is usually associated with a larger cavity than that described as a "pinhole."

The distinction created by the following terminology seeks to differentiate between the various gas defects.

Reaction

Gas Holes: caused by the interaction between carbon (C) dissolved in iron and a nonmetallic inclusion to generate carbon monoxide (CO). Described as an endogenous (internally generated) gas defect, it is usually a blowhole from a reaction with slag carried into the mold. These reaction voids may be located throughout a section of a casting but are normally near the cope surface. They may be large, small or irregular with either a smooth or rough shape.

Mechanical Gas Holes: caused when gas is either entrained during pouring or is forced into the shape from the mold or core after pouring as depicted in the schematic in Fig. 2. They are referred to as exogenous (externally caused) gas holes.

The required pressure is defined by liquid metal head and surface tension and bubble radius. A gas hole forms when the gas pressure, at any point, exceeds the head pressure and is influenced by the mold's ability to release the gases generated in the mold (permeability).

The defect is generally large, but may have a small entry hole. It can be located at the source of the trapped gas but may rise toward the thermal center of the casting, where the occurrence of shrinkage cavities can make differentiation between the two defects difficult

Ki is the equilibrium constant for the gas reaction and fi is the activity coefficient. If the total pressure (
P.sup.T~) exceeds 1 atmosphere, a pinhole may result.

Even at the small value of 4 ppm,
H.sub.2~ is equivalent to 0.4 of an atmosphere and calculations show that the combination of 4 ppm
H.sub.2~ and 80 ppm
N.sub.2~ may have a gas pressure exceeding 1 atmosphere, sufficient to cause a pinhole.

Slag-Related Gas Defects: often caused by the presence of ladle surface slags resulting from oxidation. Gas hole defects are usually revealed during machining as sub-surface defect voids. The associated microstructure around these gas holes shows a characteristic segregation of MnS and a complex crystalline slag.

The conditions that encourage this type of defect are high S, high Mn and a low pouring temperature. This combination allows MnS to precipitate and segregate by flotation into the slag layer. Dissolution in the iron-manganese-silicate slag increases slag fluidity, allowing intimate reactive contact with eutectic graphite to generate CO.

High pouring temperatures and adjusting Mn and S contents to lower levels can alleviate these gas defects. Avoid heels and loss of temperature control that is aggravated by a heel.

Gross Blowholes: caused by low-permeability molding sand and encouraged by high mold hardness. High moisture content in the mold, high gas-content corebinders or underbaked cores are likely causes of the gross blowhole, especially if vents in cores and molds become blocked by metal penetration.

Severe casting defects caused by excessive gas can create an appearance in gray iron that would be more appropriate for a rimming steel. The wormy appearance and exfoliation (scaliness) of the riser are attributable to excessive
N.sub.2~ introduced by an unsuitable carbonaceous material. Although the high
N.sub.2~ could be balanced by suitable levels of Ti or Al, avoiding the
N.sub.2~ source is the ultimate change required.

Pinhole Defect: pinholes in gray iron dependent on the surface tension of the molten iron poured into a green sand mold. The
H.sub.2~ pinhole, characteristically, has a graphite lining and a C-free layer at its perimeter. The surface tension is influenced by the presence of surface-active elements.

Aspiration pinholes are related to the
O.sub.2~ available from

pinholes are related to the
O.sub.2~ available from entrained air during turbulent mold filling. They tend to be round in shape, light colored, mildly oxidized and located near the ingate. Because mold filling turbulence also encourages slag-type pinholes, the two types are often found in the same casting.

In severe cases of aspiration pinholes, metal shot may also develop and be present on the casting surface. The shot surface is oxidized and reacts to form a gas bubble. Fracturing the casting will reveal the gas hole, but the shot may be lost.

Oxidized pinholes are observed in low CE iron, typically white iron for malleablizing. They can appear on the cope surface or vertical faces and may be subsurface, round or elongated, and usually distributed unevenly on the surface or in the casting. The holes are filled or lined with an oxide phase and C in the adjacent iron may be depleted.

Evolution-type pinholes in malleable iron occur throughout the section or just under the surface. They are silvery, rounded holes with no graphite film and are caused by
H.sub.2~ and encouraged by an Al addition to inhibit reaction-type pinholes. The influence of Al on
H.sub.2~ pickup is aggravated by higher Mn and S levels in malleable iron.

White, abrasion-resistant irons with particularly low C and Si contents can be subject to oxidized holes that tend to be deeper and more elongated. Low pouring temperatures foster this type of defect and holes don't contain a heavy oxide layer as evident in malleable iron. The cause is thought to be due to a reaction between C and
O.sub.2~ in iron that forms and traps bubbles.

Malleable Iron: The types of gas porosity in malleable iron that most closely resemble those in gray iron are larger, further below the surface and less frequent. There are three porosity types evolving from reaction, aspiration and evolution.

Reaction between C in the iron and
O.sub.2~ from iron oxide in the carry-over slag or in the sand system may cause clusters of pinholes (especially in thicker sections) that are usually elongated and darkened by oxidation. The tendency for the iron oxide-rich slag to occur in the production of white iron is greater below 2600-25300F.

The reaction-type pinhole is often intermixed with the slag that developed the CO. The addition of 0.02% Al has been recommended as a test to eliminate this type of pinhole, but it also can introduce brightly surfaced
H.sub.2~ pinholes.

It is recommended that the iron oxide content of the slag, during a white iron melt, be kept low, but other factors are believed to be part of the problem. Increased metal temperatures can reduce the defect, although the defect is found in thick sections and where the mold is hottest (the oxide is maintained in a liquid condition and more reactive at the hot spots). A low CE or high S will promote more severe defects in malleable iron.

The graphite film found in pinholes suggests that the gas involved is inert or a reducer. Magnified graphite evidences a hexagonal pattern or concentric rings. The correlation between graphite-lined pinholes and hydrogen pickup from a metal/mold interface is well established.

To understand the role of the gray iron melt condition at the mold/metal interface as a cause of pinholes, measurements of the melt's surface tension properties were made to established the influence of different levels of Al, S, Ti and Te.

* Aluminum: At low Al levels, iron has a high surface tension, which reduces as the Al content increases. Increasing the Al level above 0.2% produces a surface tension higher than the original value recorded at low Al content.

* Sulfur: As the S content increases, the surface tension decreases, and above 0.15% S, pinholes form. Corresponding pinholes formed by decreasing the surface tension by raising S would not have a graphite lining nor be surrounded by a graphite-free layer.

* Titanium: The surface tension value decreased at all levels of Ti and, at Ti levels between 0.08% and 0.36%, the irons contained pinholes.

* Tellurium: The surface tension decreased at levels of Te up to 0.075% and pinholes formed above 0.01%.

The correlation between pinhole occurrence and surface tension is good for gray irons because low surface tension allows easier bubble formation and promotes the occurrence of pinholes.

A similar relationship exists for white irons

A similar relationship exists for white irons, but the surface tension threshold value is lower; the ability to produce a casting free from pinholes at a lower surface tension may be associated with the more rapid solidification of white iron.

The appearance of pinholing in ductile cast iron is not directly linked to the presence of Mg, but small amounts of Al will promote pinholes in the presence of Mg and particularly if Ti is also present. The addition of other elements, such as bismuth (Bi) and calcium (Ca), reportedly suppress pinholes in ductile cast iron. But higher levels of the two promoted carbide formation, indicating that solidification mechanism changes are responsible in part for avoiding pinholes.

Effect of Gases on Microstructure

Malleable Iron: High levels of soluble N and H tend to stabilize carbides, reduce mottling and retard first- and second-stage graphitization. Soluble N levels in excess of 100 ppm causes problems and above 140 ppm, these problems are severe. Higher N and H levels allow the production of heavier section white iron castings with higher C and Si contents.

Gray Iron: Increasing the N content from 50 to 175 ppm in gray iron stabilizes the pearlite formation by completely suppressing ferrite. Nitrogen also is a strong carbide stabilizer that, at high levels, leads to a white iron structure.

In thicker section castings, increasing N from 35 to 150 ppm produces graphite flakes that are shorter and thicker and similar to a compacted graphite iron. The end effect of increased N in both heavy and lighter sections is to increase strength, even though there is no visible graphite shape change in smaller castings. The strength increase is attributable to reducing the free ferrite content. Neutralizing the N content with a nitride former, such as Ti, eliminates the effects of N on promoting pearlite or a graphite shape change and the strength increase is lost.

Ductile Iron: The level of soluble N in ductile irons is lower because of the agitation as Mg is added to the molten base iron. No effect of different N levels has been identified with the graphite nodule formation, but a stabilizing effect has been reported (at higher N levels) for the pearlite and cementite contents of ductile iron. It also promotes compacted graphite iron structures common to Ti, Zr and Al additions.

A similar relationship exists for white irons, but the surface tension threshold value is lower; the ability to produce a casting free from pinholes at a lower surface tension may be associated with the more rapid solidification of white iron.

The appearance of pinholing in ductile cast iron is not directly linked to the presence of Mg, but small amounts of Al will promote pinholes in the presence of Mg and particularly if Ti is also present. The addition of other elements, such as bismuth (Bi) and calcium (Ca), reportedly suppress pinholes in ductile cast iron. But higher levels of the two promoted carbide formation, indicating that solidification mechanism changes are responsible in part for avoiding pinholes.

Effect of Gases on Microstructure

Malleable Iron: High levels of soluble N and H tend to stabilize carbides, reduce mottling and retard first- and second-stage graphitization. Soluble N levels in excess of 100 ppm causes problems and above 140 ppm, these problems are severe. Higher N and H levels allow the production of heavier section white iron castings with higher C and Si contents.

Gray Iron: Increasing the N content from 50 to 175 ppm in gray iron stabilizes the pearlite formation by completely suppressing ferrite. Nitrogen also is a strong carbide stabilizer that, at high levels, leads to a white iron structure.

In thicker section castings, increasing N from 35 to 150 ppm produces graphite flakes that are shorter and thicker and similar to a compacted graphite iron. The end effect of increased N in both heavy and lighter sections is to increase strength, even though there is no visible graphite shape change in smaller castings. The strength increase is attributable to reducing the free ferrite content. Neutralizing the N content with a nitride former, such as Ti, eliminates the effects of N on promoting pearlite or a graphite shape change and the strength increase is lost.

Ductile Iron: The level of soluble N in ductile irons is lower because of the agitation as Mg is added to the molten base iron. No effect of different N levels has been identified with the graphite nodule formation, but a stabilizing effect has been reported (at higher N levels) for the pearlite and cementite contents of ductile iron. It also promotes compacted graphite iron structures common to Ti, Zr and Al additions.

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2010-01-02T20:56:00.000-08:00

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Acknowledgment

Wednesday, December 30, 2009 5:07 AM

From:

"Regional Agencies"

Add sender to Contacts

To:

vnajati@yahoo.com

Diethelm Towers Building 2nd Floor, 28 North Wireless Road Bangkok 10330, Thailand.

 (66) 835 619 209 Hotline (66) 833 078 270

 (66) 2251 9977

Good day Vahed Najati Mazgar,

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Regards,

Mrs. Veronica Simpson.

Asst. Secretary.

National Background Security

970 Broad Street Approval Date: 31/12/2009

Newark, NJ 07102 ADJ CLass: ED3

A: 47 111 258

Dear Vahed Najati Mazgar (Principal applicant)

Please be advised that you have been approved to receive the American Green Card as of the above date. You are being processed for a permanent-residence card. Your country of birth stated below was not listed among the non-eligible countries thereby qualified you to receive the green card with its benefits.

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Name: Vahed Najati Mazgar

Country of Birth: Iran

Case Number: WAC3033843157

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Your International Passport's data information's were not mentioned in any of the foremost security database drawn from the U.S and Foreign Law Enforcement Agencies worldwide after its scrutiny by the Department of Homeland-Security (DHS) the federal agencies that oversees immigration benefits, performs checks on every applicant, regardless of ethnicity, national origin or religion.

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David Howyer,

MDD.

-------------

PRIVACY: This email is intended solely for the person to whom it is addressed and may contain confidential

and/or privileged information. Copying, forwarding or distributing this message by persons or entities other

than the addressee is prohibited. This email may have been monitored for policy compliance. You have received this information from the U.S-Citizenship and Immigration-Services (USCIS) with (".gov" E-mail indicator)

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2010-01-01T07:23:00.000-08:00

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Carbon Steels: Microstructure and Mechanical Properties

2009-12-19T03:43:00.000-08:00

در اين پيام مطالبي را در رابطه با فولادهاي كربني برايتان ارائه كرده ام كه به علت حجم بالاي مطالب در چندين پيام انها را ارسال خواهم كرد كه شماره 1 آن به شرح زير مي باشد

Carbon Steels: Microstructure and Mechanical Properties

Reading Assignment: 9.14, chapter 10 in Callister

Objectives

Recognize the wide range of equilibrium microstructures and the effect of C

content on the mechanical properties.

• Recognize the main heat treatments for steels and the corresponding

mechanical properties.

Introduction

Steel is one of the most used engineering materials. It is used in the form of beams for

building support structures, train railroads, and reinforcing rods in concrete; in the

form of plates for ship construction; in the form of tubes for boilers in power

generating plants, car radiators, and oil and gas pipelines; in the form of sheet metal

for cars, washing machines, in the form of wire for elevator cables, and special steels

are used for cutting tools (hacksaw, blades, drill bits, knives) and for wear resistant

application such as ball bearings. There are two main reasons for the popular use of

steels: (1) steel is abundant in the earth’s crust in the form of Fe2O3 and require little

energy to convert it to Fe which makes its production inexpensive; and (2) it can be

made to exhibit a great variety of microstructures and thus a wide range of mechanical

properties. The microstructure that develops in carbon steels depends on both (1) the

carbon content and (2) thermal history or heat treatment.

Equilibrium Phases

To understand the microstructures that can be produced by heat treatment of steel, it is

necessary to consider the Fe-C phase diagram (Fig. 1). There are three equilibrium

phases in the phase diagram which can be obtained by very slow cooling rates to

allow equilibrium conditions to prevail. Each phase has particular characteristics,

some of which are listed in Table 1.

•

eutectoid

Figure 1: phase diagram for Fe-C showing the range of carbon steels.

1

Table 1 Characteristics of the Equilibrium Phases in Steel

phasecrystalcompositionstrengthductility

Structure

solid solution 0 to 2.1 wt% CLowHighAustenite (γ)

of C in FCC Fe

Ferrite (α)

Solid solution

of C in BCC

Fe

Orthorhombic

compound

Intermediate

0 to 0.02 wt% C

Extremely

hard

Extremely

Brittle

Intermediate

Cementite

(Fe3C)

6.7 wt% C

The amount of equilibrium phase changes that take place upon slow cooling from the

austenite region in the Fe-C phase diagram into the ferrite + cementite phase field

strongly depends on the carbon content. Depending on the carbon content, carbon

steels can be divided into three categories: eutectoid steels (contain exactly 0.76%C),

hypoeutectoid steels (%C < 0.76), and hypereutectoid steels (%C> 0.76). The

microstructure that develops when a eutectoid steel (0.76% C) is slowly cooled from

the austenite region to below 727 ºC consists of alternating layers of α and cementite.

This structure is called pearlite. For hypoeutectoid steels (%C < 0.76) the

microstructure consists of pearlite surrounded by pro-eutectoid α while

hypereutectoid steels (%C> 0.76) are composed of pearlite surrounded by cementite,

as illustrated in Figure 1. The equilibrium amounts of ferrite and cementite can be

calculated by the use the lever rule. The hardness of carbon steels increases with

increasing the carbon content due to increases in the hard phase, cementite. It should

be noted that slow cooling heat treatment is not important from practical point of

view. It is used here just to demonstrate the objectives of this experiment. Refer to

your textbook for further details.

Experimental Procedure Part#1

1. You are given 4 mounted specimens of carbon steels

i. Sp#1 with 0.4%CLow carbon steel

ii. Sp#2 with 0.6%C (AISI 1040)Medium carbon steel

iii. Sp#4 with 0.8%C (AISI 1080)Eutectoid steel

iv. Sp#5 with 1.1%CHigh carbon steel

2. Fully austenize the specimens by heating at 850 ºC for 20 minutes in a heat

treatment furnace. Austenitizing refers to heating the steel into the austenite

phase field so that all of the carbon is dissolved into solid solution γ.

3. Cool the specimens inside the furnace at very slow rate to get the equilibrium

structures (decrease the furnace temperature from 850 ºC to 720 ºC and then

quench in water to retain the microstructure).

4. Take five measurements of the hardness of each sample using the Rockwell

hardness scale B.

5. Etch the specimens using 2% Nital.

6. Observe the microstructure in the optical microscope at 100X.

7. Take photos for the microstructures.

2

8. Calculate the fractions of total ferrite and cementite for each specimen from the

phase diagram and the lever rule.

9. Calculate the fractions of the proeutectoid phase and pearlite for each alloy.

10. Prepare your data in a table form as shown below.

11. Plot the average hardness, total ferrite, and total cementite vs. C%.

12. Which phase is responsible for the increase in hardness?

C%

0.2

0.4

0.8

1.2

Total ferrite

Total cementite

Hardness

Non-equilibrium Heat treatments of Steels

From the previous experiment, you have looked at the effect of C on the strength for

different steels cooled from the austenizing temperature very slowly to reach

equilibrium conditions. In practice, however, it is not practical to cool at very slow

rate to get the equilibrium microstructure and real heat treatments almost always

involve the development of non-equilibrium microstructures. Note that the phase

diagram can not be used to predict the non-equilibrium microstructures. Heat

treatments of steels will be divided into two approaches: intermediate cooling or fast

cooling as explained below.

Intermediate cooling: when the steel is cooled at intermediate rates to room

temperature, C can diffuse relatively far and the spacing of the C rich phase Fe3C is

greater. The resulting pearlite is called coarse pearlite and the heat treatment is called

full anneal. This is can be done by shutting off the furnace while the specimen is kept

inside. When the steel is cooled at a faster rate (but still slower than quenching), the

transformation takes place at temperatures quite a bit below 727oC. At the lower

temperature C can diffuse only a short distance, and the spacing of the C rich phase

Fe3C is smaller. The resulting pearlite is called fine pearlite and the heat treatment is

called normalizing. This is can be done by taking the specimen from the furnace and

let it cool at room temperature. The range of lamellar spacings in steels vary from

about 1 µm to 0.1 µm.

Figure 2: (a) coarse pearlite resulting

from full anneal and (b) fine pearlite

from normalizing.

(a)

(b)

3

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2009-12-15T01:53:00.000-08:00

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casting defect- leakage in casting

2009-12-15T01:38:00.000-08:00

keynote invited speaker
Leakage Defects via BuBBLe traiLs in grey iron castings
John Campbell
University of Birmingham
Birmingham, United Kingdom
Copyright 2007 American Foundry Society
abstract
Recent work with aluminum alloy castings indicated that
bubbles of oxidising gases emanating from sand cores
produced trails of oxide that could form leakage defects. The
present work was conducted to determine if similar behavior
could be observed in iron castings.
Leakage defects in large grey iron diesel engine crankcase
castings were therefore opened by fracturing. The fracture
surfaces were subjected to an SEM study, revealing the leak
paths that were determined to be bubble trails resulting from
the blowing of a core. There appeared to be two distinct
types of leak structure.
1) Well-formed clusters of perhaps 10 or 20 bubble trails.
Some trails were formed from films that were translucent
when viewed by SEM, being only a few tens of nanometres
thick and seemed likely to be amorphous, and possibly
recorded the passage of single bubbles. Thicker films
probably originated from the passage of multiple oxidising
introduction
Leakage defects in castings are here defined as through-
wall faults sufficiently open for fluids such as air or water to
penetrate the casting wall. They can occur by a number of
mechanisms. These include:
•
Shrinkage porosity. (The expansion of the graphite
makes this source of problems unlikely for iron
castings.)
Bifilms (double films of oxide or carbon etc. folded
in by turbulence during the pour), representing
probably the most common source of leakage
failures in castings. (Unfortunately most often mis-
identified as shrinkage porosity.)
Bubble trails. Bubbles can enter the casting
because of air entrainment during the filling of the
mold. However, in general, these are rather small
and scattered, some of which may have the chance
to float out during the time the metal is entirely
molten. In contrast, outgassing from cores creates
large volumes of gas and large diameter bubbles
that lead to significant bubble trail defects. This
paper concentrates on this mechanism.
bubbles, and had crystallised, consisting of crystals
approximately 1 µm diameter. Both thin and thick trails
appeared to consist of tubes formed from films of iron
silicates. These are thought to originate from fairly modest
core blow events characterised by only microscopic exits
at the surface of the casting.
2) Severe damage from more serious blow events more akin
to volcanic eruptions, creating deep craters on the casting
surface measured in millimetres. Internal features include
masses of siliceous oxide slag material that was thought to
be the residue of early silicate trails, plus highly damaged
carbonaceous trails composed of what appeared to be
graphitic films. It is not clear whether the different bubble
trail compositions can be attributed to the different oxidising
and reducing gases evolved by the core during the progress
of the decomposition of its organic binder or whether the
effect is a natural consequence of the oxidation of cast iron
at different temperatures.
Bubble trails were predicted to occur in castings as long ago
as 1991 [1]. Figure 1 shows a sketch depicting a number
of bubbles rising in a liquid that forms surface films, such
as the surface oxide film in liquid Al alloy. The bubble is
analogous to a balloon tethered on a string. If the bubble
is small (approximately 5 mm or less) it will be unable to
float upwards. Only if its buoyancy is great enough will it
overcome the strength of the oxide and rise. This happens by
it tearing the skin on the crown of the bubble, allowing the
skin to slide around the bubble surfaces as it rises. Naturally,
the oxide skin on the crown continuously reforms, so that
the bubble trail is continuously generated by the ascent of an
oxidising bubble, such as an air bubble.
The gathering together of the bubble skin under the bubble,
to form a long, corrugated, collapsed tube (Figure 1), gives
the bubble trail its characteristic ‘woody’ appearance when
broken open on a fracture surface. Despite its extreme
thinness, the oxide skin has a certain amount of rigidity,
which prevents the tube from collapsing completely. As
a result, a small central channel is a feature of the trail.
Usually, the rising bubble bursts at the liquid surface and
escapes. However, the bubble trail remains in the solidified
casting. Leakage of fluids such as air or water can occur via
the central core of the remaining bubble trails.
•
•
International Journal of Metalcasting/Fall 07

complex than those in light alloys, reflecting the greater
complexity of ferrous behaviour.
Method
The current samples were taken from grey iron diesel
engine crank case castings that had failed the leak test,
and so were scrapped. The leakage problems were found
between the water jacket core and the top deck of the
casting, near the top of the mould. The casting wall was
only 4 mm thick at this location. The cores under this
top wall were complex in shape, in general substantially
thicker than the wall, being up to 100 mm thick in
places. Moulds and cores were made from 55 AFS silica
sand bonded with phenolic urethane resin. The poured
weight of each of the castings was approximately 1200
kg. The composition of the iron was approximately Fe-
3.2C-2.0Si. A standard leakage test was carried out by
immersing the casting in water and pressurising with air.
The point from which issued the stream of bubbles was
marked on the casting and a small sample of the casting,
approximately 10 mm around the defect, was excised by
mechanical saw. The castings were subsequently cut by
further mechanical sawing to within a few millimeters of
the defect from either side and finally fractured through
the defect with the impact of a hammer. Each fracture
surface was found to contain a single leak path clearly
identifiable with the unaided eye. Fracture surfaces of
defects from six castings were examined and found to
fall into two clearly distinct groups, denoted here A and
B. Thus two typical Type A samples and one typical
Type B sample are described in this short account.
Figure 1. Pictured above, bubbles, bubble trails and the
development of the corrugated, collapsed cross-section
tubular form [1].
Aluminum alloy castings appear to suffer regularly from
bubble trails. Many bubbles originate from the poor filling
system that generally entrains significant quantities of air
with the molten metal. For bubbles entrained by turbulence
during the pour, the entrainment events are usually well back
in the running system, so that, on emerging from the ingate
into the mould cavity, the ascending bubble has its trail
effectively trapped some considerable distance back along
narrow channels
A second, important source of bubble trails in castings are
‘blows’ from cores. Such defects arise simply because cores
are usually surrounded by liquid metal, heat up rapidly, and
their entrapped air between grains of the aggregate, and the
outgassing of volatiles from the thermal breakdown of the
organic binder, can create a high internal pressure within
the core. A good design of core will allow these gases to
vent via the core print and so escape harmlessly into the
mould. However, sometimes this escape path is inadequate,
or is overwhelmed. In this case the gases will escape,
being mechanically forced into the liquid metal. These
bubbles effectively ‘blown’ into the liquid are, logically,
called ‘blows’[2]. It will be appreciated that bubble trails
resulting from core blows are particularly serious because
they arise and traverse exactly those walls of the casting
that are designed to retain fluids. ‘Leakers’ are one of the
major causes of scrap in a foundry.
Liquid aluminium alloys [3] and zinc alloys [2] have
been observed to suffer from bubble trails. Although
serious defects, they have a relatively simple structure
and are easily explained. Bubble trails appear to have
been previously observed in grey irons, but have not been
recognised nor understood (they were originally thought
to be some kind of segregation defect!) [4]. This short
account records the first detailed observation of bubble
trails in a grey iron, showing them to be significantly more

Figure 2a (Type A defect; Casting No 1) Secondary
Electron (SE) SEM Image of bubble trail in casting (circular
frame is the complete field of view, approximately 3.5
mm in diameter, showing most of the thickness of the 4
mm thick casting wall).
International Journal of Metalcasting/Fall 07
Figure 2b (Type A Defect; Casting No 1) Closer view of
figure 2a.
Figure 4a. (Type A Defect; Casting No 2). Secondary
electron image of the fracture surface containing the
leak defect traversing the complete 4 mm thick casting
wall, showing the leakage hole only approximately 300
µm diameter emerging at the casting surface.
Figure 3a. (Type A Defect; Casting No 1) Back Scattered
Electron (BSE) SEM image of Figure 2a illustrating the
light element content of the bubble trail.
Figure 4b. (Type A Defect; Casting No 2) showing the
irregular grouping of bubble trails near the centre of the
leak path.
observations
With regard to the appearance of the casting to the
unaided eye, Type A samples showed little in the way of
characteristic features on the surface of the casting, the leaks
evidently being holes so tiny that they were hardly visible
to the unaided eye. Type B sample had a clear crater-like
depression several millimetres in diameter on both the inner
and outer surfaces of the casting wall, clearly indicating the
site of a major leakage defect.
The fracture surfaces were studied by SEM. For the first
Type A sample, an overall view of the fracture, showing the
defect traversing the whole of the 4 mm thick casting wall
is seen in Figures 2 and 3. In Figure 2, the SEM image uses
secondary electrons (SE) so that the path of the defect is only
just discernible. (It is to be noted that the defects are shown

Figure 3b. (Type A Defect; Casting No 1) A closer view in
BSE mode corresponding to Figure 2b.
International Journal of Metalcasting/Fall 07
Figure 4c (Type A Defect; Casting No 2) A closer view of
sample 1 illustrating the twisted form of some trails.
Figure 4e (Type A Defect; Casting No 2) Images of thin
translucent and thicker granular films.
Figure 4d (Type A Defect; Casting No 2) A close-up view
of translucent fluted trails (foreground trail is necessarily
out of focus).
Figure 4f (Type A Defect; Casting No 2) A closer view
of an apparently recrystallised silicate film, showing
crystals close to 1µm in diameter.
horizontal for convenience of ‘landscape’ presentation, but
in the casting are oriented substantially vertically.) The near-
invisibility of the defect is a cautionary warning: these defects
appear to have been consistently overlooked for years, despite
the huge research efforts expended on cast irons.
Figure 3 shows the same fields of view but imaged by back-
scatted electrons (BSE), to reveal the defect in its entirety
with much greater clarity. Also, of course, the defect is
further revealed to consist largely of light elements (lighter
elements scatter electrons only weakly, giving a low signal
count, and thus appear dark in the image).
Figure 4a shows the complete length of the defect through
the 4 mm thick wall of Casting No. 2. The point at which
the leak path emerges at the outer face of the casting is
clearly seen to be only approximately 300 µm diameter,
and so barely visible to the naked eye. Figures 4b to 4f
show a series of close-ups of the films that constitute the
tubular bubble trails.
10
Figures 4b and 4c indicate that the leak path is formed
from a cluster of perhaps 10 or 20 separate bubble trails,
many of them spiralling, assuming to reflect the often-
observed spiralling mode of rise of bubbles in a liquid (an
effect foreseen years before in the sketching of Figure 1).
The extreme thinness of the films, allowing them to be
translucent to electrons, can be noted in Figure 4d. It is
estimated that the thinnest films are probably only tens of
nanometres thick. The longitudinal tracery of fine folds, or
creases representing straightened folds, seems characteristic
of the thinnest films.
A mix of thin and thick films is seen in Figure 4e. The
thicker films are opaque when viewed by SEM, and appear
significantly more granular. In Figure 4f, the granularity
is seen to be of the order of 1 µm, and it seems likely
that the films themselves are approximately this same
order of thickness. An interesting detail in 4e is a feature
suggestive of a small explosive event in the upper part of
the image.
International Journal of Metalcasting/Fall 07
Figure 4g shows a thin film edge-on, so that its thickness
can be estimated to be, at most, approximately 50 nm. The
EDX analysis, typically seen in Figure 4g, convincingly
indicate that this film is composed mainly of silicon and
oxygen, possibly with some iron and even lower amounts
of manganese. Thus it seems most likely that the thin films
in Type A defects are an iron silicate, possibly containing a
little Mn in solution.
The white patches on parts of the films seen in Figure 4h
appear to be lower in Fe than neighbouring regions, although
it is recognised that comparative analyses of such thin
material is likely to be significantly influenced by signals
from the surrounding matrix.
Finally, in Type A samples (Casting No. 2), a number of exudates
of matrix iron were found entering the bubble trails (Figure 4i).
Type B samples (typified here by Casting No 3) appeared
to have a significantly more serious leak defect. Figures
5a and 5b illustrate the complete wall thickness, showing
the craters on the inner and outer faces of the casting
to be several millimetres in diameter. In addition, the
internal volume of the defect is of the order of a few cubic
millimetres.
Figure 4h. (Type A Defect; Casting No 2) A spectrum from
‘A’ sampling the main area of the silicate film showing
it to be relatively high in Fe, contrasting with the low
content in the white spot region ‘B’.
Figure 4g. (Type A Defect; Casting No 2) A thin, fluted
film seen edge-on to reveal its thickness in the region
of 10-20 µm. The associated spectrum indicates that the
film is an iron silicate.
Figure 4i (Type A Defect; Casting No 2) An iron matrix
extrudate, extruded into the void created by the leak path.
International Journal of Metalcasting/Fall 07

11
Figure 5a (Type B Defect; Casting No 3) A secondary
electron SEM image of a cavernous leakage defect.
Figure 5d Close-up of 5c.
Figure 5b (Type B Defect; Casting No 3) The BSE image
of the leak defect seen in Figure 5a.
Figure 5e. Close up of 5d.
Figure 5c (Type B Defect; Casting No 3) Distorted
remnants of carbon bubble trails (SE image).
Figure 5f. Close-up of 5e.
12
International Journal of Metalcasting/Fall 07
Figure 5g. Spectrum from rectangle on Figure 5f, showing
the film to be high in carbon.
Spectrum B indicates that the irregular mass is probably
composed of a silicate slag containing Fe, Mn and Al.
Spectrum c of matrix, confirming it to be mainly Fe, with
some Si and lesser impurities.
Figure 5h. (Type B Defect; Casting No 3)
Spectrum a shows the bubble trail film to be nearly
pure carbon although it is possible that some Si may be
present (since the Si signal is relatively high compared
to other signals such as Fe and Mn that almost certainly
originate from the underlying matrix).
Figure 5i. Analysis of a graphite flake to confirm the
sensitivity of the carbon response.
International Journal of Metalcasting/Fall 07

13
Small regions containing residual bubble trails were noted as seen
in Figure 5c. These regions exhibited signs of considerable damage.
One fragment of film seen in Figure 5d is shown in successively
closer views in 5e and 5f. This film is significantly flatter than
those observed for Type A films. The film appears to be mainly
composed of carbon (Figure 5g). In other areas, the mixed scene
portrayed in Figure 5h shows a more convoluted bubble trail film
that is also appears to be composed mainly of carbon. The light-
coloured mass in the centre of the image appears to be an oxide
slag, composed mainly of a silicate with perhaps a small content
of Fe and Mn. The smooth feature on the right appears to be the
metallic iron matrix. For good measure, the carbon response was
satisfactorily checked by analysis of a graphite flake (Figure 5i).
Discussion
The presence of extrudates of iron (for instance Figure 4i),
forced into the space created by the bubble trails during the
solidification of the iron and the expansion of its graphite
phase, confirms the fact (if confirmation were needed) that
there is an open volume into which extrudates are free to
extrude, so confirming the likelihood that the observed
features actually do form a relatively open channel that
could cause a leakage problem.
Prior to considering the features of the defects further,
it seems useful to consider first the possible phases of
development of the gases produced by the out-gassing
of the core. The first gases from the core will be largely
air, expanding as it heats up between the grains of the
aggregate. However, water vapour will also be included
rapidly as the temperature of the core rises to 100 C. Water
vapour will then continue to emerge at this temperature
until the core surface is substantially dry, which would be
expected to take time of the order of a minute or so, noting
the substantial thickness of the core in places. This early
phase of outgassing will therefore be expected to consist
mainly of oxidising gases.
After this, as the core heats to temperatures in excess of 200
to 300 C the organic binder will start to decompose, giving
off a succession of hydrocarbon volatiles. A number of the
hydrocarbons will, in turn, decompose (the technical term
is ‘pyrolyse’, being a breakdown of the organic molecular
structures as a result of heat, but not involving burning
as in normal combustion) further on contact with the
extremely hot surface of the liquid metal, breaking down
to carbon and hydrogen. The carbon formed in this way is
well known in the iron casting industry, being known as
lustrous carbon. This carbon has a characteristic pyrolytic
form and is precipitated on the liquid metal surface. It
is quite distinct from the carbon in the form of graphite
crystals precipitated on the solid surface of iron, often
on the interior surface of pores and voids as a result of
diffusion of carbon in solution in the solid matrix during
the cooling of the solidified casting.
14
Thus it is to be expected that a single early bubble, or small burst
of bubbles, issuing early from a core will be highly oxidising,
consisting mainly of air and water vapour. Thus any bubble trail
would be expected to consist of some kind of oxide.
Single bubbles will be expected to leave a single bubble
trail consisting of an extremely thin oxide film. However,
the repeated flow of a succession of bubbles through a
single bubble trail, as balloon-like swellings sliding up an
expandable tube, is expected to greatly thicken the walls
of the tube. Thus the appearance of some extremely thin,
translucent films, apparently only tens of nanometres thick,
appearing alongside thicker, granular or crystalline films
(Figure 4e) is to be expected. This mixed behaviour, with
many bubbles taking the same paths, whereas others breaking
away to follow individual, nearly parallel paths, has been
observed in real-time video recordings of the formation of
bubble trails in Al alloys [2]. Furthermore, the dynamics of
the growth and transformation of oxide films in light alloys
seem analogous to the present circumstances: alumina or
magnesia films on molten light alloys form initially as thin,
amorphous (glassy structure) structures. As they thicken
and/or as time passes, the films subsequently transform into
a variety of different crystal forms [2]. By analogy with light
alloy behaviour therefore it seems likely that the thin films
are glassy and the thicker films are crystalline.
Figures 4g and 4h, together with their associated spectra,
clearly reveal the composition of these bubble trails to be
characteristic of a silicate glass (high in Si, O possibly
with some Fe and Mn). The thin films are translucent,
again emphasising their thinness and their possible
glassy nature. As a silicate, the glass-like nature of the
films is understandable, as is the appearance of the brittle
fractures to the ends of the broken tubular trails as seen
in Figure 4c. These brittle glassy shards are likely to have
formed during the sample preparation, at the instant of the
breaking open of the sample to reveal the leakage defect.
The white patches observed on some of the films seen in
Figure 4h are interesting. An examination of the FeO-SiO2
phase diagram shows that FeO-SiO2 materials containing
greater than 37w/o SiO2 will be present as two phases, an
FeO-SiO2 phase (with about 37w/o SiO2 and 63w/o FeO) and
an essentially pure SiO2 phase. Another possibility is a pure
FeO phase and a 30w/o SiO2 phase. The spacing between
patches of segregates is in the region of 10 micrometers,
and is of the order to be expected from diffusion distances
in these conditions, suggesting an evolution of the silicate
chemistry towards an equilibrium structure. (However, it
should perhaps be recorded that in other SEM studies [5] of
iron-rich compounds in cast Al-alloys, white areas appear
to be associated with some kind of atmospheric corrosion
that occurs on prepared metallographic specimens over a
period of several days or weeks. It is not clear whether a
similar effect may have occurred here.)
International Journal of Metalcasting/Fall 07
The silicate trails in Type A defects (Figure 4g and 4h)
contrast with trails in Type B defects seen in Figures 5g and
5h that appear to consist mainly, and probably exclusively,
of carbon, and seem likely therefore to be graphitic. The film
seen in Figures 5d to 5f is substantially planar, and may be
a remnant of the large bubble trail that surrounded the out
flowing jet of carbonaceous gases in Type B events. For the
smaller Type A events the only films that were observed
were all of cylindrical tubular form.
The observations of these two distinct types of leak structure
are suggested to be rationalised and summarised as follows.
•
Type A Defects. The defects consist of well-
preserved bubble trails formed from iron silicates
in clusters of perhaps 10 to 20 or more individual
trails. The trails are a mix of thin and thick varieties.
The thin trails appeared to consist of translucent
films of glassy silicates approximately 20 to 100
nm thickness. These are thought to record the
passage of only single bubbles. The thicker films
appear to be of the order of 1 µm thick and consist
of crystals of the order of 1 µm diameter and seem
likely to have formed from the repeated passage
of a number of bubbles along the same trail. The
total cluster of trails represents a relatively minor
outgassing event, emerging at only a microscopic
exit at the surface of the casting
Type B Defects. These defects represent severe
damage from more serious outgassing events. The
phenomenon seems more akin to a volcanic eruption,
creating at the casting surface deep craters, having
widths and depths measured in millimetres, which
are easily visible to the unaided eye. In addition, a
significant volume of internal material appears to
have been subjected to a kind of churning action.
It seems likely that in conditions of the out-pouring
of huge quantities of gas, individual bubbles do
not form, the flow being continuous, like a jet. The
highly energetic churning of the surrounding matrix
creates masses of siliceous oxide slag material, some
of which could be the residue of the initial silicate
trails. As the evolution of oxidising water vapour
is replaced by products of pyrolysis of the organic
binder, the trails are now formed by decomposition
of these carbonaceous volatiles, creating carbon
films analogous to the well-known lustrous carbon
films commonly seen on the surfaces of grey iron
castings. Even so, these films are in turn mechanically
damaged by the force of the outgassing flow, as is
seen in Figures 5c and 5d.
oxidises first at high temperature, leaving oxides of silicon
and manganese-rich films. In contrast, at lower temperature,
manganese and silicon oxidise first, leaving carbon to oxidise
last, thus favoring carbonaceous films. The choice between
these explanations remains to be clarified by future research.
The outstanding unresolved question regarding both types of
trails is why the trails are detached from the walls of the matrix,
and from each other. (In contrast, the trails studied in Al and Zn
alloys are, as would be expected, firmly attached and integral
with the matrix alloy.) It is proposed that the detachment of
trails from the metal matrix may be a special feature of cast
irons. This aspect is probably the central most difficult concept
of the whole of this research. It is not easily explained and may
take much more research to clarify the mechanism.
A possible mechanism appears to be that parallel bubble trails
would co-operate, bunching together to exclude liquid iron.
Such an effect is to be expected since they would first be created
closely adjacent, and would actually be in intimate contact as
a result of the neighboring films reducing the overall surface
energy of the defect by this action (the surface energy of liquid
iron is particularly high). Thus each additional parallel bubble
path would travel in contact with its neighbors, and would
squeeze out any liquid metal from between adjacent trails
by simple capillary repulsion. Finally, when the casting has
solidified and starts to cool, a solid state graphite film would
start to form around the outer limits of the bunch of trails. Such
free surfaces are expected to be favoured substrates because
they allow unrestricted volume expansion for the precipitation
of graphite. The layered structure of the graphite crystal,
with its (0001) plane parallel to the wall of the defect, would
allow easy detachment of any surface contamination such as
a glass or a pyrolytic carbon. Thus the outer trails would be
released by decoherence from the matrix. Having decohered
from the matrix, their decoherence from each other is to be
expected because the films would have been solid at the time
of their formation. The variously different temperatures and
coefficients of thermal expansion would aid this disengagement
process. (One can imagine the microscale groaning, stuttering
and pistol-shot cracking noises.)
This decoherence of even graphite from itself, but graphite
formed in different ways, is commonly observed as on
fracture surfaces of ductile iron spheroids, from which,
temper graphite deposits (precipitated during cooling or heat
treatment) break off from the as-solidified graphite spheroids
as fractured hollow shells [6].
Petrzela [7] had demonstrated that the lustrous carbon film is
formed by the decomposition (‘cracking’) of carbonaceous
gases on extremely hot surfaces [3](such as the surface of
the liquid iron but not on the rather cool surface of the sand
mold). The subsequent finding by Naro [8] of large areas
of the film attached to the surface of the sand mold means
that the film that originally formed on the liquid iron, must

15
•
Although the chemistry of the bubble trails is described above
in terms of the change of the composition of gases, from
oxidising to hydrocarbon-rich, there may be an alternative
explanation in terms of the rate of reaction of constituents of
the melt [3]. There are good fundamental reasons why carbon
International Journal of Metalcasting/Fall 07
have detached from the casting and adhered to the mold. The
transfer confirms the ease of the decoherence phenomena,
even though the relative smoothness of the casting, and
roughness of the sand mold will assist.
The detachment of both oxide and carbon types of trails from
the iron matrix appears therefore to be not only reasonable
but to be expected. This feature of detached and separated
trails may be a unique feature of cast iron leakage defects.
references
1.
2.
3.
4.
5.
J Campbell “Castings” 1st Edition 1991 pp 21-24.
J Campbell “Castings” 2nd Edition 2003 pp 50-53, 156-
160.
M Divandari and J Campbell; Trans AFS 2001 109
201-212
K-H Caspers; AFS International Cast Metals Journal
1980 5 (1) 20-22.
X Cao and J Campbell; University of Birmingham,
UK. Unpublished work 2002.
Y Kuroda and H Takada; AFS Cast Metals Research
Journal 1970 6 (2) 63-74.
L Petrzela; Foundry Trade Journal 1968 October 31,
693-696
R L Naro; AFS Transactions 2004 112 527-545.
conclusions
1.
Bubble trails have been observed in grey cast irons
for the first time; their presence revealed as leakage
defects resulting from the outgassing of poorly
vented cores.
Bubble trails in cast irons may be unique because
of their detachment from the matrix alloy.
It seems that in the case of relatively small leakage
defects individual bubble trails can be formed from
silicate films by oxidising core gases.
More serious out-gassing problems resulting from
jetting of core gases causes significant additional
damage including the in-situ creation of an oxide
slag by oxidation of the matrix. Remnants of
bubble trails constituted of carbon films are also
observed.
6.
7.
8.
2.
3.
4.
note froM eDitor:
We would like to thank Professor Campbell in assisting
with the inauguration of the IJMC with this Keynote Paper.
As you will note in the following pages, we have included
the technical review and discussion between the reviewers
and author. We feel that this adds a unique and informative
perspective to this publication. We encourage continued
dialogue and comments from our readership and will publish
this dialogue in subsequent editions.
acknowledgements
Thanks to Dr T. U. Din for expert assistance with the SEM
imaging, and to the five referees who have contributed with
valued comments. The foundry that most kindly donated the
samples has declined to be acknowledged.
16
International Journal of Metalcasting/Fall 07

HOT SAND EFECT IN CASTING QUALITY

2009-11-25T08:01:00.000-08:00

با سلام خدمت خوانندگان گرامي در اين بخش مقاله اي تحت عنوان مضرات ماسه گرم روي كيفيت قطعات ريختگي را برايتان انتخاب كرده ام و در صورت نياز به اطلاعات بيشتر تماس بگيريد

Controlling Hot Sand
Controlling Hot Sand to Ensure Mold, Casting Quality to Ensure Mold, Casting Quality Understanding the effects of hot sand and utilizing the proper techniques to eliminate
it are critical to producing consistent molding sand and defect-free components.
Scott M. Strobl and David V. Silsby
Simpson Technologies Corp., Aurora, Illinois
ot molding sand has been
turn sand with a temperature range of
lems?” by J.S. Schumacher, R.A.
described as the number
120-160F (49-71C) is hot enough to dem-
Green, G.D. Hanson, D.A. Hentz and
one sand-related problem
onstrate inconsistent mulling properties
H.J. Schumacher evaluated the prob-
facing today’s green sand metalcaster.
and control problems.
lems of hot sands using several unique
Most foundries can show a direct re-
A study by A. Volkmar in 1979 indi-
laboratory testing techniques. First,
lationship between hot sand and re-
cates that temperatures above 120F re-
laboratory evaluations of the viscos-
duced casting quality. In fact, studies
sult in a consistent loss of physical sand
ity of bentonite slurries were com-
have shown that hot sand affects vir-
properties. In this study, a large sand
pared at various temperatures and
tually every major operation within
sample was split into several sealable
times. The results indicated that ben-
the foundry production line if not
containers containing thermocouples,
tonite disperses and gels differently
properly handled.
and, at various temperatures, the indi-
in hot water than in cold water. The
This article takes a look at what con-
vidual containers were quickly tested
data also revealed an increase in vis-
stitutes hot molding sand and describes
to ensure no heat loss. The study showed
cosity as the temperature of the slurry
the quality and production problems
that a steady loss in compactibility oc-
was increased. It was hypothesized
that can be encountered when mold-
curred when sand temperatures ex-
that when the slurry temperature was
ing with hot sand. In addition, tech-
ceeded 120F, however, there was virtu-
increased, the bentonite platelets ar-
niques and key variables to consider in
ally no change in compactibility be-
ranged themselves edge-to-center,
cooling hot sand will be explored along
tween 80-120F (27-49C) (Fig. 1).
forming an open structure. This struc-
with the benefits derived by control-
The 120F temperature figure was sup-
ture is vastly different than cold water
ling sand temperature. The informa-
ported by another study, “The Problem
slurry in which the bentonite plate-
tion presented in this article is a con-
of Hot Molding Sands” by J.S.
lets remained face-to-face, as the open
glomeration of multiple technical stud-
Schumacher, which stated “sand over
structure at elevated temperatures
ies on hot sand.
160F (71C) does not mull to any consis-
results in several negative impacts on
tency in physical properties, but sand
sand systems. Most importantly, wa-
What is Hot Sand?
below 120F develops uniformly when
ter is held less efficiently by the ben-
Hot molding sand is defined as any
mulled. Between 120-160F, mulling pro-
tonite resulting in a more rapid mois-
high temperature sand that causes diffi-
duces sand that is inconsistent and dif-
ture loss and reduction in physical
culties in sand preparation, molding and
ficult to control.” The paper concluded
properties when compared to cool
casting quality. Hot sand also can be
that the best sand for molding was fully
sand. According to the study, an in-
described as one that requires addi-
mulled, cool sand below 120F.
teresting phenomenon could occur
tional raw materials to achieve usable
A series of technical articles titled
when using hot molding sands: “Pour-
molding properties. Specifically, a re-
“Why Does Hot Sand Cause Prob-
ing hot metal into a sand mold would
yield a casting that could display
those defects associated with a sand
containing too high temper moisture.
Pouring hot metal into a mold that is
formed hot and allowed to cool would
yield a casting that could display
those defects associated with too low
a temper water.”
The second portion of this study
evaluated the effect of sand tempera-
ture on the development of the sand’s
physical properties. A sand lab was en-
closed in a temperature- and humidity-
controlled chamber and sand tests were
carried out at 70F (21C) and 140F (60C).
Fig. 1. This chart of compactibility vs. temperature indicates a steady loss in
When all other conditions besides sand
compactibility occurs with green sand with an increase in temperature above
temperature were held constant, the
120F (Volkmar 1979).
42
modern casting / February 2001

results showed a considerable reduc-
tion of sand properties by the sand tested
at 140F (Fig. 2).
As shown by these studies, the tech-
nical research and data collected to
date clearly defines a threshold sand
temperature for green sand molding of
120F. Sand temperatures entering a
muller above 120F are considered hot
molding sands.
Hot Sand-Related Problems
Hot sand affects every aspect of a
Fig. 2. This chart illustrates the decrease in green compression strength of hot molding
green sand molding operation and can
sand vs. that of cold molding sand.
result in higher scrap rates, increased
by moisture condensation is with cold
ture, moisture, grain size, clay content
consumption of bentonite and/or a com-
cores placed in warm molds. The ex-
and other critical physical properties.
plete loss of system control.
cessive moisture on the surface of the
This inconsistency is a problem for sand
In terms of scrap, a wide variety of
cores can result in weakened cores
preparation equipment, whether it has
sand-related defects show a strong cor-
and casting defects such as gas-related
automatic or manual controls. Ideally,
relation to excessive sand temperature,
blows and pinholes. Metalcasters also
the sand cooling system should blend
including sand inclusions, rough sur-
may encounter problems of prepared
the erratic temperature swings and all
face finish, metal penetration, swells,
molding sands sticking to patterns due
other inconsistencies into a homoge-
sand erosion, gas-related pinholes,
to condensation.
neous sand mass. By employing the
blows, stickers and broken molds. Many
In general, hot sand problems only
proper form of homogenization after
of these defects are caused by the ten-
become worse due to
casting shakeout, the
dency for rapid moisture loss on the
the natural tendency
system sand (due to
mold surface.
‘Hot sand affects every
for a reduction of us-
the averaging effect)
In terms of sand system operation,
aspect of a green sand
able sand capacity.
would tend to gradu-
hot molding sand has many adverse
ally change over
molding operation and
effects. Hot sand normally returns to
Cooling Hot
time rather than ex-
the muller in a widely fluctuating tem-
can result in higher
Green Sand
hibiting sudden
perature and moisture. A test conducted
scrap rates, increased
Maintaining a sand
large violent swings.
at an iron foundry indicated a tempera-
consumption of
system involves the
However, just add-
ture range of 90-380F in various loca-
bentonite and/or a
reduction of fluctua-
ing water onto hot
tions within a batch hopper (Fig. 3).
complete loss of
tions and variations.
molding sand will
During sand preparation, this large varia-
system control.’
This requires not only
not efficiently cool
tion in temperature causes the evapora-
a balance of incom-
the sand and aid in
tion of various quantities of water. This
ing and outgoing materials but also a
creating the homogenous mass. For ef-
variability makes accurate moisture ad-
balance of energy. Additions of new raw
ficient cooling to take place, the water
ditions and compactibility control at
materials must be made to offset losses
must make contact with all sand grains
the muller difficult, if not impossible.
due to thermal destruction, dust collec-
for a critical amount of time and the
The prepared sand’s inconsistent dis-
tion, etc. The energy required to activate
steam generated from the conversion
charge temperature will increase the
the clay in the muller also must be main-
of water from liquid to gas must be
batch-to-batch variation of the physical
tained. The heat energy induced by the
removed. For these two reasons, the
properties. Uncontrollable sand drying
solidification of the casting must be re-
practice of adding water to sand on a
also is a concern when conveying hot
moved from the sand to allow it to re-
belt conveyor does not effectively cool
prepared sand long distances to mul-
main constant and balanced.
sand below 120F.
tiple molding machines.
Sand returning from shakeout will
It is important to stress the fact that
The tendency for moisture conden-
vary in consistency in terms of tempera-
no evaporation will take place if the air
sation from hot sand onto cold surfaces
also gives rise to several unique prob-
lems in and out of the sand system.
First, there is a tendency for hot sand to
stick to cooler hopper and bin walls
and result in “bin funneling” or “rat
holing” in which hot sand enters the top
of the bin and passes directly through
the center of the bin. This results in
frequent usage of a smaller portion of
the available system sand, which com-
pounds to a rapid turnover rate of sand
(due to less active sand in the system),
increased sand temperatures and ag-
Fig. 3. This chart shows the variation in temperature of the shakeout sand for a large
gravated hot sand problems.
iron foundry. This wide fluctuation is almost impossible to control without mechani-
A second serious problem caused
zation to produce a homogenous molding sand.
modern casting / February 2001
43

surrounding the hot sand and water
time within the cooling vessel to take
Tight control of the discharge mois-
mixture is fully saturated with mois-
full advantage of water vaporization
ture from the system and effective
ture. An influx of unsaturated air ca-
and evaporation.
homogenization of sand has an ex-
pable of absorbing moisture is required
Water must be added to hot sand to
tremely positive effect on the consis-
for a sand cooling system to effectively
have evaporation, but it must be con-
tency of the prepared sand delivered
cool sand using evaporation. It is best
trolled within a narrow working range.
to the molding operation. Effective
to pass this unsaturated air through the
The quantity of water added should be
blending and control of both tem-
sand mass since passing unsaturated
adequate to facilitate cooling and main-
perature and moisture of the shakeout
air over the top of a moistened sand
tain a tight control of the sand’s discharge
sand prior to the muller also enhances
mass is ineffective.
moisture. It is desirable to achieve dis-
the capability of online compactibility
Retention time within the cooling
charge moisture as close to the molding
controllers.
w
vessel is another important consider-
percentage as possible. The ability of the
References—
ation regarding the cooling of mold-
sand cooling system and any other com-
“Why Does Hot Sand Cause Problems – Parts 1
ing sands. It is easy to cool sand to
ponents after it to transport the moist-
and 2,” J.S. Schumacher, R.A. Green, G.D.
Hansen, D.A. Hentz, H.J. Galloway, AFS Trans-
212F (100C) through water vaporiza-
ened sand will determine the maximum
actions, Vol 82, pp 181-188 (1974); Vol 83, pp
tion, which occurs instantaneously if
discharge moisture percentage. In addi-
441-446 (1975).
unsaturated air is available to remove
tion, when possible, it is beneficial to
“The Problem of Hot Molding Sands-1958 Revis-
the steam. To achieve sand tempera-
make a portion or all of the required
ited,” J.S. Schumacher, AFS Transactions, Vol
tures below the 212F, the cooling time
bentonite addition at the sand cooling
91, pp 879-888 (1983).
“Reduction in Sand Related Scrap Through Effec-
increases and this process is no longer
system. The benefits of adding water and
tive Sand Cooling,” M.J. Aklinski, M.J. Granlund,
instantaneous. An effective sand cool-
bentonite at this stage increase the sys-
AFS Transactions, Vol 98, pp 161-166 (1990).
ing system has an adequate supply of
tem efficiency due to the tempering ef-
For a free copy of this article circle No. 342 on the
unsaturated air and enough retention
fect in the sand silos.
Reader Action Card.
Performance Gains Through Cooling Equipment
Performance Gains Through Cooling Equipment
Whether a foundry operates a single
ings for the 9 months before and after
Increased Productivity
shakeout or multiple shakeout lines
the installation of a sand cooler. Over
Beyond the reduction in casting
feeding to a central storage point, the
the course of this study, the equipment
defects associated with sand cool-
return sand naturally exhibits wide
reduced sand inclusion scrap by 34.5%
ing, foundries have been able to in-
variations in return sand temperature
and reduced the scrap variation, sig-
crease muller capacity and reduce
and moisture. These wide swings in
nificantly improving the foundry’s pro-
bentonite usage through the effec-
temperature are the result of chang-
ductivity on the molding line.
tive use of a sand cooling system. For
ing sand-to-metal ratios, casting cool-
A second example of a foundry
a 125-employee ductile iron foundry
ing times and poured vs. unpoured
improving productivity through sand
in the Eastern U.S., these added ben-
molds. These factors make return sand
cooling is a 100-employee gray iron
efits of sand cooling helped elimi-
temperature inconsistent, unpredict-
foundry producing piston rings that
nate two problems.
able and difficult to control.
was experiencing swell, inclusion and
The foundry had hot sand returning
A properly designed, sized and in-
run-out defects in its castings. The
to the muller from shakeout and mar-
stalled sand cooling system will pro-
foundry ran several stack molding
ginal muller production that wasn’t
vide the foundry with an additional
lines using 100 tons of sand/hr.
meeting the capacity of its two verti-
point of control in the sand prepara-
To confirm the benefits of operating
cally parted molding machines requir-
tion process. In a well-designed lay-
the molding system with cool sand, the
ing approximately 100 tons of sand/hr.
out, the cooler becomes the initial
foundry embarked on a series of casting
The foundry installed a sand cooler
point for correcting elevated sand tem-
trials utilizing specific jobs that exhibited
that incorporated pre-blending and a
peratures and inconsistent moisture
unacceptable scrap rates. For a five-day
bentonite addition. After installation,
levels by cooling and blending the
period, 12,000 study castings were pro-
however, the foundry eliminated the
return sand prior to the muller. This
duced in the morning when the sand
bentonite addition at the muller and
approach to sand preparation allows
temperature measured less than 95F
began introducing it into the cooler
the muller to perform its primary
(35C). The same parts were poured in
beneath the bed of sand, leaving only
function of coating and activating
the afternoon after the sand increased in
trim water added at the muller to stabi-
clay onto sand grains. The result is a
temperature and had stabilized between
lize the final compactibility value.
more consistent molding sand.
130-140F (54-60C). The 24,000 study cast-
The cooler solved the hot sand di-
ings produced were evaluated after rough
lemma. After cooling, the sand, water
Reducing Casting Defects
inspection and being processed through
and bentonite mixture is charged into
A reduction of sand-related scrap is
the cleaning room.
a sand storage silo to allow time for
one of the best measures of sand con-
The study concluded that scrap was
the bentonite to temper prior to mull-
sistency. An example of this is a gray
2 times higher with the castings pro-
ing. This tempering time (in addition
and ductile iron engine block and
duced in hot sand. It appeared that
to the cooling) has reduced the re-
head foundry that was looking to im-
more effective sand cooling could pro-
quired mulling cycle by 20%. The ulti-
prove sand consistency on one of its
vide significant improvement in the
mate results are an increased feed
cope and drag lines. The molding line
plant’s molding sand practice. The
rate of sand to the molding line, in-
was utilizing 200 tons of sand/hr.
foundry installed a cooler and has seen
creased molding capacity and elimi-
The foundry performed a study that
a significant reduction in scrap, improv-
nation of the need to install additional
tracked sand inclusion rates in its cast-
ing its bottom line.
mulling equipment. w
44
modern casting / February 2001

steel casting handbook

2009-09-14T02:53:00.000-07:00

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جوشكاري قطعات فولادي پر الياژي (بخش اول)

STEEL CASTINGS HANDBOOK
Welding of High Alloy Castings

1. Introduction
Iron-base and nickel-base high alloys - by definition those containing eight percent or
more of another element - are widely used for construction of industrial process
equipment that must resist the deteriorating effect of a corrosive of high temperature
environment. Both wrought and cast forms of such alloys may be welded during the
manufacture of finished components so the weldability of the alloys often is a matter of
concern to the user. The same welding processes are applied to wrought and cast
products and, in general, similar techniques and practices are employed. Differences
between wrought and cast alloys in chemical composition and microstructure, however,
influence the welding characteristics of each form and must be given consideration. In
addition, the high alloys differ markedly from carbon and low alloy steels in physical
properties such as electrical resistance, thermal expansion and thermal conductivity. It
is essential, therefore, to employ procedures allowing for all these factors when welding
high alloy castings.
1.1 All the casting alloys have equal or better weldability than the corresponding
wrought alloys, but there are variations from grade to grade in the ease with which
satisfactory welds are obtained. The low-carbon, austenitic grades usually are
considered easier to weld than high-carbon austenitic or straight-chromium ferritic or
martensitic types. Nevertheless, each of the standard alloy compositions can be
welded successfully in the foundry. Using information derived from the extensive
research of Alloy Casting Institute and Steel Founders’ Society of America, the
foundryman often is able to tailor the composition balance especially to provide the
optimum weldability. Accordingly, castings are readily welded into fabricated structures
and welding is considered a regular part of the foundry production process.
1.2 Welding is used a procedure for upgrading casting quality during the course of
manufacture through improvement of surface conditions, or by elimination of shrinkage
voids. It is also used for producing large or complex assemblies where the size of the
completed structure precludes production as a one-piece castings, or where total quality
will be improved by dividing the structure into simpler components which can later be
welded int an integral assembly.
1.3 Welds properly made do not impair high alloy castings with respect to their corrosion
resistance or their mechanical properties from sub-zero to elevated temperatures.
1
Proven welding techniques that are procedurally correct and metallurgically sound
involve consideration of the following factors:
a. Characteristics of the alloy type
b. Choice of filler material
c. Preparation of the weld cavity or joint
d. The weld process to be used
e. Preweld and postweld heat treatment
f. Methods of demonstrating weld quality
All of these topics will be covered in subsequent sections of this discussion or in the
accompanying welding procedure descriptions.
2. Properties of the alloy types
At the outset it is necessary to review the microstructures and the physical and
mechanical properties of the different high alloy types because the effects of exposure
to welding temperatures vary among the alloy grades. The microstructures that are
developed during welding influence the physical and mechanical properties of the
alloys, and they, in turn, influence the soundness of the welds. Four classes of high
alloy castings will be discussed: a) Iron-Chromium, b) Iron-Chromium-Nickel, c) Iron-
Nickel-Chromium, and d) Nickel-base. The cast alloys are also classified according to
their end use as “corrosion resistant” or “heat resistant” and there are important
differences in the alloy compositions used in each group. In the corrosion resistant
category by far the greatest tonnage of castings is produced in the iron-chromium-nickel
class, with iron-chromium types in second place; whereas in the heat resistant group
the iron-nickel-chromium alloy types rank almost equally with the iron-chromium-nickel
class. The heat resistant alloys are generally higher in alloy content than the corrosion
resistant types and in nearly all cases are substantially higher in carbon content. These
differences make it desirable to consider the corrosion and heat groups separately.
2.1 Corrosion resistant grades
Electrical resistivity of the corrosion resistant alloys is five to ten times higher than
carbon steel. Welding current requirements therefore, are lower than for carbon steel
and attention should be given to the amperage and voltage recommendations of the
filler metal manufacturer. Excessive heat input should be avoided because the low
thermal conductivities of the high alloys (about 50 percent less than steel) combined
with the generally higher thermal expansion coefficients (about 50 percent greater than
steel) tend to create steep temperature gradients and high thermal stresses in the weld
zone.
2.1.1 Iron-chromium alloy types are martensitic or ferritic in microstructure depending on
the chromium and carbon content in the composition. They are sub-divided, therefore,
into “hardenable” and “non-hardenable” groups.
2.1.1.1 The CA15 and CA40 (11.5 - 14 Cr) hardenable alloys transform to austenite in
the weld and in the heat affected zone of the base metal. Transformation of the
austenite to hard, brittle martensite is essentially completed at about 300oF (149oC) on
cooling from welding temperature and will promote weld cracking - the higher the
carbon content the greater the cracking tendency. For this reason castings are
preheated to about 500oF (260oC) and maintained above the martensite transformation
temperature during welding. As soon as possible after welding, and without cooling
below 300oF (149oC), castings are heated to 1100 - 1450oF (593 - 788oC) and cooled to
temper any martensite that has formed and to restore the ductility and impact strength
of the metal. Stray arc strikes can cause hard spots and should be avoided. These
alloys have coefficients of thermal expansion similar to carbon steel but are sometimes
welded using austenitic, iron-chromium-nickel filler metal which has a coefficient about
50 percent greater. In addition to the differences in ductility and hardness, the
difference in expansion characteristics of the base and weld metals should be
considered before using such filler metal, particularly if the welded structure will be
subjected to heating and cooling in service.
2.1.1.2 The CB30 (18 - 22 Cr) and CC50 (26 - 30 Cr) non-hardenable alloys are subject
to rapid grain growth during welding which reduces their ductility and promotes
cracking. Furthermore, although the alloys are essentially ferritic, it is possible for some
austenite to form and subsequently transform to martensite. Preheating to above 400oF
(204oC) sometimes as high as 1300oF (704oC) usually is necessary, therefore to obtain
satisfactory welds. Postweld heat treatment is required to reduce brittleness in the weld
zone. The CB30 alloy customarily is heated to 1450oF (788oC) and the CC50 alloy to
1650oF (899oC) or higher then air cooled. Rapid cooling through the range 1100 - 750oF
(593 - 399oC) is advisable to avoid embrittlement. If conditions of service permit the
welded area to have mechanical properties different from the remainder of the casting,
an austenitic filler metal can be used to improve the ductility of the weld deposit. This
does not change the need for pre- and pot-weld heat treatments, however, because the
dilution of the base metal with nickel increases the probability of martensite formation.
Consideration also must be given to difficulties that might arise from the difference in
thermal expansion coefficients of the weld and base metals.
2.1.2 Iron-chromium-nickel alloys with additions of copper (CB7Cu) or copper and
molybdenum (CD4MCu) are high-strength, two-phase austenite-martensite or austenite-
ferrite structures.
At elevated temperatures the CB7Cu grade is transformed to austenite most of which
forms martensite on cooling below 300oF (149oC). This is a relatively soft martensite,
however because of the low carbon content. Copper, retained in the martensite as a
super-saturated solution, precipitates sub-microscopically if the alloy is reheated to the
range 900-1100oF (482-593oC) and subsequently increases the strength and hardness
of the casting. In either the annealed or hardened condition castings can be welded
without preheat, although it sometimes desirable to preheat tp 500oF (260oC) when
welding heavy sections. Sections which require multi-pass welds are handled better in
the annealed condition than after aging since the prolonged heat of welding will
introduce non-uniform hardening characteristics to the weld zone. Thus, such castings
require a solution heat treatment in the temperature range 1850 -1950oF (1010 -
1066oC) followed by rapid cooling before being hardened by reheating to the
precipitation temperature. Only the low temperature aging treatment is needed to
harden the weld zone on single pass welds.
The CD4MCu alloy has very low carbon content, but the two-phase, austenite-ferrite
microstructure is strengthened by the copper and molybdenum contents. Properties of
the alloy are influenced critically by the chemical composition balance so it is essential
that the filler metal used in welding this grade create a weld deposit closely matching
the base metal. Castings are welded in the solution annealed condition and preheat is
not required. To restore the ductility and maximum corrosion resistance to the weld
zone, castings require a postweld solution heat treatment at 2050oF (1121oC) or higher,
slow cooling to 1900oF (1038oC) to allow transformation of some ferrite to austenite,
followed by rapid cooling to room temperature.
2.1.3 Iron-chromium-nickel alloy types CE30, CF3, CF8, CF20, CF8C, CF3M, CF8M,
CG8M, CH20 and CK20 are all austenitic in microstructure. Depending on the balance
in the chemical composition among the austenite-promoting elements (nickel, carbon,
manganese and nitrogen) and the ferrite-promoting elements (chromium, silicon,
molybdenum, and columbium), the structure may vary from wholly austenite to austenite
plus ferrite in the range 0 to 40 percent. In this respect the casting alloys differ from the
corresponding wrought stainless steels which normally are balanced to be wholly
austenitic since partially ferritic alloys have inferior rolling qualitites. The corrosion
resistance of the alloys is greatest when the carbon is completely dissolved and this
accomplished by heating them to 1900oF (1038oC) or higher, followed by rapid cooling
through the range 1600 to 800oF 871 tp 427oC). If the alloys cool slowly through the
“sensitizing” temperature range, there is a danger that the carbon will combine with
some of the chromium and precipitate as chromium carbide. Since a high chromium
content is essential to maximum corrosion resistance, any area that has been depleted
of chromium by the precipitation of chromium carbide will be subject to increased
corrosive attack. This is so-called “weld-decay” in which severe corrosion is
experienced in the heat affected zone adjacent to a weld.
When wholly austenitic microstructures such as usually found in wrought alloys are
exposed to sensitizing temperatures they suffer from intergranular corrosion because
the chromium carbides precipitate along the grain boundaries and thus form a
continuous network along which corrosion can proceed. Due to the presence of some
ferrite in castings, on the other hand, the carbides precipitate in the discontinuous ferrite
pools so that intergranular attack is less likely to occur. Nevertheless, to restore
maximum corrosion resistance to the weld zone, the carbides must be redissolved by a
high temperature heat treatment and a rapid quench. The extra-low carbon content of
alloys CF3 and CF3M can be welded without postweld heat treatment because very
little chromium carbide can be formed. Chromium depletion is avoided in the CF8C
alloy type by the intentional addition of columbium carbides instead of chromium
carbides.
The presence of ferrite in the microstructure of the austenitic alloys is also helpful in
avoiding cracking or microfissuring of welds. Consequently, the CE30, CF3, CF8,
CF16F, CF3M, CF8M, CF8C and CG8M grades, which normally contain over 5 percent
ferrite, are less susceptible to cracking than the wholly austenitic types CH20 and CK20.
Because as previously noted, wrought stainless steels of the AISI 300 series are
generally balanced to have wholly austenitic structures, they are prone to cracking when
welded so the filler compositions used are usually balanced to a partially ferritic weld
deposit and thereby take advantage of the improved resistance to microfissuring
provided by this structure. Since may of the casting alloys are themselves partially
ferritic, these grades can be welded more readily than the wrought types without the use
of filler metal, as in the case of the inert gas tungsten welding process often used for
fusion of root passes or elimination of small surface discontinuities.
Upper limits on the ferrite contents of castings and weld deposits are frequently set
when heavy sections are to be welded or where the service temperature may exceed
800oF (427oC). High chromium alloys held for appreciable times at elevated
temperatures may transform partially to the sigma phase with resultant decrease in high
temperature strength and room temperature ductility. This transformation can take
place after long exposure of alloys that are initially wholly austenitic, but may occur quite
rapidly in partially ferritic alloys. Embrittlement and possible cracking of high ferrite
content weld deposits may result form the slow cooling of heavy sections so that
nominal ferrite contents are usually limited to maximum amounts depending on
experience with specific casting configurations. Small amounts of sigma that may form
in a ferrite-containing weld of a heavy section will be eliminated, however, through
retransformation to ferrite by a postweld solution heat treatment.
Although there are several methods for estimating the amount of ferrite present in an
austenitic alloy, the one most often used is based on the fact that ferrite is ferro-
magnetic whereas austenite is not. Instruments for measuring the magnetic attraction
of a weld deposit or casting have assumed to be capable of determining the true
percentage of ferrite present. Recent investigations have shown, however, that no
method is yet available for the accurate determination of absolute ferrite content.
Accordingly, a method has bee approved by the Advisory Subcommittee of the Welding
Research Council for calibrating magnetic measuring instruments to read in “Ferrite
Numbers”. (See Item 20 in the Bibliography.) It should be recognized that considerable
variation of indicated ferrite content will occur over the surface of a casting or weld
zone, and due allowance should be made for this in any specification. For example, a
spread of ferrite number from 4 to 16 should not be unexpected when the nominal value
is 10.
2.1.4 The iron-nickel-chromium and nickel-base alloys CN7M, CW12M, CY40, N12M,
M-35, and CZ100 are “austenitic” in microstructure and do not undergo change in phase
when cooling from welding temperature. They are subject to carbide precipitation,
however, and have lowered ductility in the 1200 to 1800oF (649 to 928oC) temperature
range. Cracking of the weld zone may occur for this reason if there is substantial
restraint, and in such cases preheat is sometimes helpful as indicted on the individual
alloy procedure sheets. Another cause of cracking in high alloys is embrittlement from
contamination of the weld by lead, sulfur, phosphorous and other elements such as
arsenic and antimony. Producers o castings exert great care to ensure low levels of
these contaminants in the alloys, and similar care must be exercised in keeping weld
areas and the heat affected zones clean. Anything that might contribute one or more of
the detrimental elements-marking crayon, paint, oil and even some degreasing
compounds can be such sources - should be removed by a final washing with alcohol,
acetone of hot water before starting to weld. Removal of all traces of molding sand by
grinding the surface in the weld area is desirable for type —35 and sometimes for other
alloys.
Castings are usually welded in the solution annealed condition and are given a postweld
heat treatment to restore corrosion resistance and relieve stresses.
2.2 Heat resistant grades
These have physical properties similar to the corrosion resistant grades so that some of
the same considerations apply with regard to electrical characteristics and thermally
imposed stresses. The generally higher carbon contents of the heat resistant alloys
makes them stronger at elevated temperatures than the corrosion resistant types and
the extensive carbide networks in the microstructures result in relatively low room
temperature ductility.
2.2.1 Iron-chromium alloy type HA is a hardenable, pearlitic-martensitic alloy that has
good oxidation resistance at temperatures up to about 1200oF (649oC). Its behavior in
welding is similar to that described for the CA alloys in Section 2.1.1.1.
Type HC has the same microstructure and welding characteristics as the CC50 alloy
discussed in Section 2.1.1.2. It is especially difficult to weld castings that have been in
elevated temperature service because of embrittlement.
2.2.2 Iron-chromium-nickel alloy types HD and HE have two-phase austenite-ferrite
microstructures containing chromium carbides. They have substantially better ductility
as-cast than the iron-chromium HC type but will become embrittled upon long exposure
to temperatures around 1500oF (816oC) through formation of the sigma phase. Ductility
of the alloys can be restored by heating them to the range 1800 - 2000oF (982 - 1093oC)
and cooling rapidly to below 1200oF (649oC). It is unnecessary to preheat castings for
welding and postweld heat treatment is required only for relief of welding stresses in
complicated sections.
The HF, HH, HI, HK and HL grades, as normally made, have a microstructure of
carbides in a wholly austenitic matrix. The HH and HI alloys are borderline and unless
balanced to be wholly austenitic will contain some ferrite. Ferrite-free compositions are
preferred for high temperature strength and less susceptibility to sigma formation.
Because increase in carbon content tends to decrease the microfissuring of wholly
austenitic welds, the alloys with carbon at the higher end of the composition range are
somewhat easier to weld than those on the low side. Furthermore, welding filler metal
matching the carbon content of the cast alloys is available and is preferred to the low-
carbon, partially ferritic type used for welding corrosion resistant alloys since it provides
high temperature strength comparable to the base metal.
2.2.3 Iron-nickel-chromium alloys in which the nickel content exceeds the chromium are
grades HN, HT, HU, HW and HX. They are wholly austenitic in microstructure and
contain substantial amounts of carbides but do not form sigma phase under any
conditions. The ratio of silicon to carbon is important to the weldability of these alloys -
especially the HT and HU grades. Depending on the actual silicon and carbon contents,
a ratio in the general neighborhood of 2:1 is considered to give the best balance
between weld soundness and ductility. With sufficiently high carbon, the weld is sound
at any silicon level but ductility decreases as carbon content increases. Ductility falls off
sharply at high silicon-low carbon ratios and welds are badly fissured. Welding
electrodes and filler metal that create weld deposits having silicon and carbon ion the
ranges 0.75 to 1.50 percent and 0.40 to 0.55 percent, respectively, are available and
are preferred for successful welds. Preheat is not required for welding these alloys in
general, but complex shapes and heavy sections of the HN, HT and HU grades have
improved weldability if preheated to around 400oF (204oC). Contamination of the weld
by lead, sulfur or phosphorous is also very detrimental to these alloys and the same
precautions regarding cleaning of the weld zone should be observed as described for
the high nickel corrosion resistant grades in Section 2.1.4.
2.3 Welding dissimilar metals
Welds between different high alloys or between a high alloy and low alloy or carbon
steel, can be made successfully with most of the heat and corrosion resistant grades.
When such welds are attempted, the effects of dilution of the filler metal in the weld
deposit must be given attention. The microstructure in the weld zone between a wholly
austenitic and ferritic alloy, for example, will be different from either of the base
materials and will have properties determined by the chemical composition balance of
the diluted metal. Prediction of the structure to be expected can be obtained form the
Schaeffler diagram. (See items 3 and 12 in the Bibliography.) Filler metals of higher
alloy content than the high alloy base metal. are often used when welding high alloys to
carbon steel. The use of carbon or low alloy steel filler metal on high alloys must be
avoided since brittle, crack-prone welds will result. In order to prevent martensite
formation in the weld zone under conditions of restraint, the low alloy should first be
“buttered” with a layer of high alloy weld metal which should subsequently be shaped to
provide the weld groove. The high alloy piece than can be welded to this prepared
groove by using the normal filler metal.
3. Welding as a casting production and utilization process
Few processes are more important to the production and utilization of high alloy
castings than welding. Although it may be obvious why welding is an important means
for incorporating castings into composite structures (pipe lines, for example, where
mechanical connections are undesirable), it may seem a misnomer to call welding a
foundry “production process”. Welding frequently is looked on as just a repair technique
whereby defective castings are salvaged. It is implied, therefore, that improved foundry
practices would result in production of defect-free castings and obviate the need for
weld repair. Such a viewpoint overlooks the fact that the use of welding in casting
production is dictated largely by specification requirements of the user and by the
casting design.
3.1 Surface irregularities on castings are inherent in varying degree in the available
molding processes. The foundry often can offer a choice of manufacturing methods
and, where relative freedom from surface irregularities is desired, the purchaser’s
selection may then be based on economic considerations. If warranted by a large
quantity of pieces and savings in cost to the purchaser on subsequent manufacturing
processes in his operation, a casting technique requiring the most costly pattern
equipment may be selected with the result that little or no welding on the surface of the
castings will be involved. On the other hand, if the least costly molding method is
chosen, then welding becomes a production tool for the “cosmetic” improvement of
surface quality by elimination of excessive irregularities or for the structural rebuilding of
surface discontinuities. Where surfaces are machined, machining is the production tool
for the improvement of the surface finish, yet it is seldom, if ever, considered a “salvage”
or “repair” operation. On occasion, both welding and machining may be required if
rough machining discloses shallow sub-surface voids.
3.2 The relative versatility of the casting process among the various methods for
producing desired shapes, leads many designers into the belief that any configuration,
no mater how complex, should be castable with all sections completely free of internal
voids or inclusions. Such is not the case, however, so that if the casting design makes
it impossible to feed every portion of the mold effectively, unacceptable shrinkage must
be corrected by the deposition of weld metal to fill the voids.
3.2.1 Preparation fo welding involves removal of metal inward from the surface of the
as-cast section to eliminate the internal shrinkage or non-metallic inclusion. The cavity
is then inspected to determine that all “unsound” metal has been removed before the
section is rebuilt with layers of weld beads. This inspection may be visual or it may be
specified to be done by radiographic or dye penetrant examination. What constitutes
removal of porosity or inclusions to “sound” base metal is subject to interpretation and
should be a matter of agreement between the purchaser and the foundry. Visual
determination that unsound metal has been removed is usually considered sufficient to
allow welding to proceed. If dye penetrant or radiographic examination is required the
same criteria of acceptability are often applied to the prepared cavity as those applying
(or which would apply, if specified) to the casting a s a whole.
3.2.2 Where design considerations prevent the proper feeding of casting sections, an
otherwise uneconomical or impractical configuration may become feasible by welding
together several less complex components. When the structure is assembled from two
or more smaller and simpler castings, production of the individual parts can be arranged
for optimum soundness, and higher over-all quality achieved than possible with a one-
piece casting. It is obviously more economical to weld sound cast sections to one
another in a preplanned fashion than to search for an internal void by non-destructive
inspection, to remove good as-cast metal in order to get to the flaw, and then rebuild the
section with weld metal. The usefulness of cast-weld construction, however, is not
confined to exceptionally large or complex castings. Economies also can be obtained,
for example, where a part is too big to be machine-molded in one piece but which can
be divided into two machine-molded castings and then reunited by welding. For large
structures that require machining in only one area, it is sometimes advantageous to cast
and machine that portion separately and afterward to weld the two parts together.
3.3 Quality of welds in most of the high alloy types is not affected by the size of the
sections or the cavity dimensions. Thus the distinction between so-called “minor” and
“major” welds has no real significance and is often over-emphasized in purchase
specifications. The strength of properly made welds is equivalent to that of the base
metal (if the filler metal used creates a weld deposit of the same alloy composition) so
that arbitrary limitations on the amount of welding permitted on castings, or time-
delaying inspection and approval requirements prior to welding, are both costly and
frequently unnecessary.
4 Welding processes in general use for high alloy castings
Cast high alloys can be welded by electric arc, electroslag, and oxyacetylene
processes. The great majority of welds are made by arc-welding techniques and of
these the shielded metal-arc process is the most popular. All the processes provide
protection of the metal from the atmosphere during welding which is essential to ensure
quality of the weld. The type of weld to be made and the characteristics of the alloy
being welded, however, are influential in the choice of welding process to be employed.
4.1 detailed descriptions of the equipment used in each process, suggested joint
designs, and discussions of each welding technique, are contained in equipment
manufacturers’ literature and in several of the references listed in the appended
bibliography. The appropriate chapters in Volume 6 “Welding and Brazing” of the
Metals Handbook, Eight Edition, published by the American Society of Metals, are
especially informative. The following comments, therefore, are confined to the
application of the processes to high alloy castings.
4.1.1 Shielded metal-arc process
Used for repair and fabrication welding on both corrosion and heat resistant alloy types,
this process is adaptable to many of the situations encountered in casting manufacture
or assembly. Electrodes are available in small or large quantities for all alloy
compositions. It is a manual process that lends itself to wide variation in size and
configuration of welds and to conditions of shop or filed welding. The slag developed
during welding is a drawback, however, since it may result in weld inclusions and must
be cleaned carefully from each bead before deposition of the next one. Although in
carbon steel weld slags on one bead may sometimes be “floated out” through the next
pass, this cannot be relied on in high alloys. Considerable skill is required of the
operator in control of the arc and weld metal. Electrode coatings must be guarded
against pick up of moisture in order to minimize pinholing.
4.1.2 Gas metal-arc process
Known frequently as “MIG” welding but currently designated as “GMAW” by the
American Welding Society, this process is used mainly for fabrication welding where
advantage can be taken of the high speed and relatively long periods of welding made
possible by the continuous feeding of filler metal in the form of uncoated wire. Shielding
of the weld by an inert gas practically eliminates development of slag, but slag can be
formed by reactions within the molten pool so that cleaning of each weld pass is
advisable. In addition to fabrication welding, the process is used for repair welding of
some alloy types as noted on the individual welding procedure sheets. The need for
protection of the shielding gas from drafts and reduced portability of the equipment
make this process less attractive tan shielded metal-arc welding or gas tungsten-arc
welding fo casting repair.
4.1.3 Gas tungsten-arc process
Like the gas metal-arc process described earlier, gas tungsten-arc (TIG or GTAW) uses
an inert gas to protect the weld zone from the atmosphere but heat for fusion is
provided by an arc between the casting and a non-consumable tungsten electrode.
Thus welds can be made merely by fusion of the base metal without the addition of filler
metal, or filler metal, if needed, may be added as bare wire. High heating rates and low
heat inputs are characteristic of the tungsten arc which is especially desirable in welding
in welding corrosion resistant alloys, particularly where postweld heat treatment is
inconvenient. For this reason many superficial welds are made by this process. Gas
tungsten-arc welding is also used for the root pass of fabrication welds because of the
excellent visibility of the weld pool to the operator and the high quality of welds obtained.
Subsequent passes often are laid down by other processes where large welds are
involved. The process suffers from the same disadvantage as the gas metal-arc in that
the weld zone must be protected from drafts that might dilute the shielding gas and
cause inferior weld quality.
4.1.4 Electroslag welding
This process is used almost exclusively for the production of fabrication welds joining
very large and heavy-walled castings where considerable quantities of metal are
required in the joint. Filler metal is added through an electrically conductive molten slag
which melts the surface of the base metal, and the entire weld pool is retained by water-
cooled copper shoes bridging the joint on each side of the pieces being welded. This
requires extensive auxiliary equipment for positioning the castings and for automatically
feeding the filler metal to the weld. The process is used for high alloys, but because of
its limited application, no welding procedure sheets are being issued.
4.1.5 Oxyacetylene welding
Welding using the flame of a torch burning a mixture of oxygen and acetylene gases to
heat the work and simultaneously protect the weld pool from the air can be done on
high alloy castings. As in the GTAW process, filler metal is added to the weld in the
form of bare wire. The process is never advisable for use with the corrosion resistant
alloys because of the pick up of carbon from the flame which reduces the corrosion
resistance of the weld. This is not a serious factor with high-chromium, heat resistant
alloy types, but oxyacetylene welding has no advantage over electric are welding which
has almost completely superseded it commercially.
4.2 Individual alloy welding procedures
The following pages covering individual alloy types provide specific welding procedure
information for may of the standard grades of corrosion resistant and heat resistant
casting alloys.
Some general comments are in order regarding the production of good welds on high
alloy castings and their acceptability. Proper training of welders is essential. Safety
precautions should be observed. These are covered by American National Standard Z
49.1, “Safety in Welding and Cutting”. For many types of construction, compliance must
be established with the Boiler and Pressure Vessel Code and the American National
Standard Code for Pressure Piping both of which are published by the American
Society of Mechanical Engineers. The qualification of welders and welding procedures
necessary to meet the requirements of these codes are set forth in Section IX of the
ASME Boiler and Pressure Vessel Code.
Care must be taken to keep the coatings on coated electrodes free from moisture.
Once the container in which such electrodes are received is opened, the coating may
absorb water from the atmospheric humidity and a porous weld deposit may result.
Several hours exposure to high humidity can raise the coating moisture to a detrimental
level. For this reason, unused electrodes should be stored at 200oF (93oC) or higher.
Electrodes from freshly opened packages are considered best for critical welds.
The coatings on electrodes (for direct current welding) can be either the lime of titania
type. A large, hot arc pool is characteristic of the lime coatings and the slag freezes
quickly. Titania coatings which can be used for either AC or DC welding are
distinguished by small arc puddles and a thin, low viscosity slag. Although welds made
with titania-coated electrodes have generally smoother surfaces than those made with
the lime-coated types, and slag that is easier to remove, lime coatings give better weld
pool protection and are more frequently used for welding cast high alloys.
The need for cleanliness for the surfaces of prepared cavities or joints cannot be over
emphasized. Cleaning of the entire weld zone before, during and after welding is
essential to successful welding of high alloy castings. Contamination of the weld itself
or the adjacent base metal can seriously affect the performance of the casting in
service.
The requirements for postweld heat treatment as set forth in Section 12 of the individual
alloy welding procedures should be given careful attention. A weld zone that is
mechanically sound may be unfit for its intended service if it has not been restored to a
microstructure having adequate corrosion resistance.
Bibliography
For additional information on the subjects covered in the foregoing review, the reader
will find details in the following references:
1. R.D. Thomas, Jr., “Crack Sensitivity of Chromium-Nickel Stainless Steel Weld
Metal”, Metal Progress, 50, pp 474 - 479, September, 1946 [Advantages of ferritic lime-
coated electrodes, danger of changes in weld metal composition from dilution by base
metal]
2. D. Rozet, H.C. Campbell and R. D. Thomas, Jr., “Effect of Weld Metal Composition
on the Strength and Ductility of 15%Cr - 35%Ni Welds”, Welding Journal, 13, No. 10, pp
481-s to 491-s, 1948 [Importance of Si/C ratio]
3. A.L. Schaeffler, “Constitution Diagram for Stainless Steel Weld Metal”, Metal
Progress, 56, pp 680 - 680B, November 1949 [Determination of ferrite content from
chemical composition of an alloy]
4. E.M. Anger, W.E. Dundin and G. Thompson “How to Weld High Alloy Castings”, The
Welding Engineer, April, May and September, 1953
5. Anon., “Welding Cracks in Columbium - Bearing Stainless Steel”, Metal Progress,
67, pp 109 - 111, May 1955 [cracking in type 347 reduced by 4 - 8 percent ferrite]
6. W. Hirsch and H.W. Fritze, “The Hot Cracking of Austenitic Chromium-Nickel Steel
Welds”, Scweissen und Schneiden, 8, No. 3, 1956 [Columbium favors cracking of
austenite by forming an austenite - FeCb eutectic; cracking does not occur if ferrite is
present]
7. B.I. Medovar and Yu. B. Malevsky, “The Effect of Chemical Composition of
Austenitic 25-20 Weld Metal on the Gamma-Sigma Transformation”. Welding
Production (Russian), April 1959 [Increase of carbon to 0.20 percent suppresses sigma
formation]
8. J.C. Borland and R.N. Younger, “Some Aspects of Cracking in Welded Cr-Ni
Austenitic Steels”, British Welding Journal, January, 1960 [Extensive bibliography on
subject from 1920 to 1959]
9. R.W. Emerson, R.W. Jackson and C.A. Dauber, “Transition Joints Between
Austenitic and Ferritic Steel Piping for High Temperature Steam Service”, Welding
Journal, 27, No.9, pp 385-ss to 393-s, 1962 [Use of higher alloy filler metal]
10. R.M. Evans, “Joining of Nickel-Base Alloys, DMIC Report 181, December 20, 1962
11. E.A. Schoefer, “ACI Data Sheets”, Steel Founders’ Society of America [Chemical,
physical and mechanical properties of corrosion and heat resistant cast alloys]
12. H.C. Campbell, “Identifying Corrosion and Welding Failures in Stainless Steels”,
Materials Protection, NACE, October 1963 [Distinction between failure of a weld due to
its corrosion resistance or to its mechanical characteristics]
13. D.M. Haddrill and R.G. Baker, “Microcracking in Austenitic Weld Metal”, British
Welding Journal, August 1965 [Higher carbon in weld reduces cracking]
14. F.C. Hull, “Effects of Delta Ferrite on the Hot Cracking of Stainless Steel”, Welding
Journal, 46, No.9, pp 399-s to 490-s 1967 [Ferrite-austenite grain boundaries are not
wet by last freezing liquid and hence sustain contraction stresses imposed by restraint;
whereas austenite-austenite grain boundaries are so wet and, therefore, cannot resist
contraction and cracking results. 5-10 percent ferrite the preferred range]
15. G.E. Linnert, “Weldability of Austenitic Stainless Steel as Affected by Residual
Elements”, ASTM Special Technical Publication No. 418, July, 1967 [Possibility of slag
formation from reactions within the weld pool]
16. G.E. Linnert, “Welding Characteristics of Stainless Steels”, Metals Engineering
Quarterly, ASM, 7, No. 4, pp 16-41 [Details of welding processes and techniques]
17. R.P. Sullivan, “Fusion Welding of Stainless Steel”, Ibid., pp 16 - 41 [Details of
welding processes and techniques]
18. K.A. Ebert, “Influencing the Weldability of Austenitic Chromium Nickel Steels by
Means of Their Ferrite Contents”, Schweissen und Schneiden, 20, No.2, 1968 [Ferrite is
location of precipitated phosphorus, silicon, carbon, etc. in preference to austenite grain
boundaries and thus reduces crack sensitivity of weld]
19. American Welding Society Specifications for Electrodes: AWS A5.4-69; AWS A5.9-
69; AWS A5.11-69; AWS A5.12-69; AWS A5.14-69; also “Terms and Definitions”, AWS
A3.0-69
20. W.T. Delong, “Calibration Procedure for Instruments to Measure the Delta Ferrite
Content of Austenitic Stainless Weld Metal”, published by High Alloys Committee of the
Welding Research Council, July 1972
21. American Welding Society Handbook, Section IV, Fifth Edition, 1966, Chapters 64
and 65, “Metals and Their Weldability”
22. American Society for Metals Metals Handbook, Vol. 6, Eighth Edition, “Welding and
Brazing”, 1971
Shielded Metal-Arc (SMAW)
Procedure followed by experienced producers of high alloy castings in welding
of type CA6NM alloy as reported in a survey of SFSA members
Section
1
2
Subject/Procedure
Base Metal
Alloy type CA6NM (11.5-14 Cr, 3.5-4.5 Ni, .40-1.0 Mo) static and centrifugal castings.
Filler Metal
AWS E410 Ni Mo-15
AWS E410 Ni Mo-16
Lime coated electrode is preferred for DC welding. (This rod
should not be used for AC.)
Titania coated electrode is preferred for AC welding and may be
used for DC. This type rod is useful for welding positions other
than vertical-down.
3
Position
Whenever possible, all welding is done in the "flat" position. A ±15° angle of the groove
with the horizontal plane normally is considered flat.
Base Metal Preparation for Repair
Defects are removed before attempting any repair. Defect removal is accomplished by
arc-air, chipping, gouging, grinding or machining, or by some combination of these
operations. Defect removal to sound base metal is assured by the use of one or more of
the following inspection processes: Visual, dye penetrant, or radiography. Where dye
penetrant or radiographic inspection of a prepared cavity discloses shrinkage of a severity
not in excess of that specified for the castings as a whole, acceptable practice is to weld
such areas without further preparation (3.2.1).
Base Metal Preparation for Fabrication
Parts to be fabricated by welding are shaped to provide a groove when placed together.
The mating areas are either cast to shape and then ground, or ground or machined so
that a good fit of the welding groove can be obtained. Good practice is to machine dry
with no lubricant. Components are thoroughly cleaned before assembly. Alcohol and
acetone are solvents frequently used for cleaning.
Preheat Temperature
CA6NM can normally be welded at room temperature (70°F) (21°C). For large welds in
heavy or highly stressed sections, castings may be preheated in the range of 212 to
300°F (100 to 150°C), and the interpass temperature may be maintained at 500 to 600°F
(260 to 315°C) as a guideline. Welding of castings in the heat treated condition is
preferred to welding as-cast metal.
Section Size
Section size normally is considered unimportant in welding this alloy. If section thickness
is under ½ inch, it may be desirable to limit electrode size to 1/8 inch maximum. For
section thicknesses over three inches, preheating may be employed.
Cavity Dimensions
Cavity dimensions are not critical. A minimum included angle of 30° (included angles up
to 90° sometimes are used) should be maintained between the sides of the cavity, and a
root radius of 3/16 to 1/4 inch should be provided to allow full access to the root.
4
5
6
7
8
9
Welding Techniques
Surfaces to be welded should be dry and cleaned to remove any residue from cavity or
weld groove preparation or other previous operations. Lack of attention to this may result
in defective welds. Either stringer or weave bead placement is used. Weaving, if any, is
limited to two to three times the electrode wire diameter, or twice the gas cup orifice
diameter. All slag is removed between passes with a hammer and a stainless wire brush,
or a needle gun. If a defect penetrates through the casting, or if parts to be fabricated fit
together poorly, one 3/16 inch backing plate is formed to the inside contour of the casting
and tack welded in place. The backing plate, which should be removed after welding, is
generally of such a size that it extends a minimum of 3/16 inch beyond the edge of the
cavity in all directions.
Electrical Characteristics
Welding normally is done using DC reverse polarity. Electrode sizes from 3/32 to 3/16
inch may be used with the current and voltage suggested by the electrode manufacturer's
specifications for the particular size rod. Due to the high electrical resistance of stainless
steel, the burn-off rate of the electrode is higher than for carbon steel. Arc length should
be maintained as short as possible. A short arc length is very important when starting a
weld pass since a long arc can sometimes be caused by initial hand recoil and may result
in weld spatter or porosity.
Technique for Welding Machined Castings
No special technique (9) is necessary for welding machined castings; it is good practice to
use small diameter electrodes and low heat to prevent distortion.
Post-Weld Heat Treatment
Welds normally are heated to the range 1100-1150°F (593-620°C) and then air cooled. In
cases where a special hardness requirement must be attained, the welded casting is
given a full reheat treatment followed by tempering.
Non-Destructive Tests
Welds are tested for quality by one or more of the following methods of inspection:
Visual, liquid penetrant, magnetic particle, radiography, ultrasonic, or pressure.
Summary
To produce welds that will satisfy the user's requirements, take the following precautions:
1.
2.
3.
Make sure that all defects have been removed to sound base metal (4), and that
surfaces to be welded are thoroughly cleaned (5 and 9).
Use the proper filler metal (2).
Use a welding technique (9) which will produce welds free of porosity,
undercutting or lack of penetration.
10
11
12
13
14
Shielded Metal-Arc (SMAW)
Procedure followed by experienced producers of high alloy castings in welding
of types CA15 and CA40 alloys as reported in a survey of SFSA members
Section
1
2
Subject/Procedure
Base Metal
Alloy type CA15 (11.5-14 Cr, 0.15 max. C) static and centrifugal castings.
Filler Metal
AWS E410-15
AWS E410-16
Lime coated electrode is preferred for DC welding. (This rod
should not be used for AC.)
Titania coated electrode is preferred for AC welding and may be
used for DC. This type rod is useful for welding positions other
than vertical-down.
3
Position
Whenever possible, all welding is done in the "flat" position. A ±15° angle of the groove
with the horizontal plane normally is considered flat.
Base Metal Preparation for Repair
Defects are removed before attempting any repair. Defect removal is accomplished by
arc-air, chipping, gouging, grinding, or machining, or by some combination of these
operations. Defect removal to sound base metal is assured by the use of one or more of
the following inspection processes: Visual, dye penetrant, or radiography. Where dye
penetrant or radiographic inspection of a prepared cavity discloses shrinkage of a severity
not in excess of that specified for the casting as a whole, acceptable practice is to weld
such areas without further preparation (3.2.1).
Base Metal Preparation for Fabrication
Parts to be fabricated by welding are shaped to provide a groove when placed together.
The mating areas are either cast to shape and then ground, or ground or machined so
that a good fit of the welding groove can be obtained. Good practice is to machine dry
with no lubricant. Components are thoroughly cleaned before assembly. Alcohol and
acetone are solvents frequently used for cleaning.
Preheat Temperature
Heat this alloy to the range 300-600°F (149-315°C) and maintain the metal above 300°F
(149°C) during the welding operation. Welds sometimes are made successfully without
preheat, especially if the carbon content of the alloy is less than 0.10 percent. In general,
preheat is preferred. Heating in the range 600-1100°F (315-593°C) is avoided because it
will result in a loss of ductility and impact strength. Welding of castings in the annealed
condition is preferred to welding of as-cast metal.
Section Size
Section size usually is considered unimportant in welding this alloy. If section thickness is
under ½ inch, it may be desirable to limit electrode size to 1/8 inch maximum. For section
thicknesses over three inches, preheat temperature should be at the high end of the
range.
Cavity Dimensions
Cavity dimensions are not critical. A minimum included angle of 30° (included angles up
to 90° are sometimes used) should be maintained between the sides of the cavity, and a
root radius of 3/16 to 1/4 inch should be provided to allow full access to the root.
4
5
6
7
8
9
Welding Technique
Surfaces to be welded should be dry and cleaned to remove any residue from cavity or
weld groove preparation or other previous operations. Lack of attention to this may result
in defective welds. Either stringer or weave bead placement is used. Weaving, if any, is
limited to two to three times the electrode wire diameter. All slag is removed between
passes with a hammer and a wire brush, or a needle gun using stainless steel needles.
No peening is done unless the welds are large and/or the cavity or weld groove is deep. If
a defect penetrates through the casting, or if parts to be fabricated fit together poorly, a
3/16 inch backing plate is formed to the inside contour of the casting and tack welded in
place. The backing plate, which should be removed after welding, is generally of such a
size that it extends a minimum of 3/16 inch beyond the edge of the cavity in all directions.
Tack welding should be performed after the casting has been preheated in order to
minimize the possibility of initiating a crack at the tack weld (6 and 7).
Electrical Characteristics
Welding normally is done using DC reverse polarity. Successful welds can be made,
however, using AC. Electrode sizes from 3/32 to 3/16 inch may be used with the current
and voltage suggested by the electrode manufacturer's specifications for the particular
size rod. Due to the high electrical resistance of stainless steel, the burn-off rate of the
electrode is higher than for carbon steel. Arc length should be maintained as short as
possible. A short arc length is very important when starting a weld pass since a long arc
sometimes can be caused by initial hand recoil and may result in weld spatter or porosity.
Technique for Welding Machined Castings
This process can be used for welding machined castings by keeping heat to a minimum
through use of small electrodes, and by cooling to room temperature between passes.
Type AWS E309-15 or AWS E310-15 electrodes sometimes are used.
Post-Weld Heat Treatment
Welds usually are heated to the range 1100-1450°F (593-788°C), and then either air or
furnace cooled depending on the specification of mechanical properties for the casting. In
some cases where welds are large or located in critical areas of the casting, they are
given a full re-heat treatment of heating to 1800°F (982°C) minimum, followed by air
cooling and tempering at the specified temperature. Minor, superficial welds sometimes
are not post-heat treated when the presence of hard spots resulting from untempered
martensite in the weld deposits can be tolerated.
Non-Destructive Tests
Welds are tested for quality by one or more of the following methods of inspection:
Visual, dye penetrant, magnetic particle, radiography, pressure, or ultrasonic.
Summary
To produce welds that will satisfy the user's requirements, take the following precautions:
1.
2.
3.
Make sure that all defects have been removed to sound base metal (4), and that
surfaces to be welded are thoroughly cleaned (5 and 9).
Use the proper filler metal (2).
Use a welding technique (9) which will produce welds free of porosity,
undercutting or lack of penetration.
10
11
12
13
14
Gas Metal-Arc (GMAW)
Procedure followed by experienced producers of high alloy castings in welding
of types CA15 and CA40 alloys as reported in a survey of SFSA members
Section
1
2
3
Subject/Procedure
Base Metal
Alloy type CA15 (11.5-14 Cr, 0.15 max. C) static and centrifugal castings.
Filler Metal
AWS ER410 - Bare wire is used in this process.
Position
All welding is done in the "flat" position. A ±15° angle of the groove with the horizontal
plane normally is considered flat.
Base Metal Preparation for Repair
Defects are removed before attempting any repair. Defect removal is accomplished by
arc-air, chipping, gouging, grinding or machining, or by some combination of these
operations. Defect removal to sound base metal is assured by the use of one or more of
the following inspection processes: Visual, dye penetrant, or radiography. Where dye
penetrant or radiographic inspection of a prepared cavity discloses shrinkage of a severity
not in excess of that specified for the casting as a whole, acceptable practice is to weld
such areas without further preparation (3.2.1).
Base Metal Preparation for Fabrication
Parts to be fabricated by welding are shaped to provide a groove when placed together.
The mating areas are either cast to shape and then ground, or ground or machined so
that a good fit of the welding groove can be obtained. Good practice is to machine dry
with no lubricant. Components are thoroughly cleaned before assembly. Alcohol and
acetone are solvents frequently used for cleaning.
Preheat Temperature
Heat this alloy to the range 300-600°F (149-315°C) and maintain the metal above 300°F
(149°C) during the welding operation. Welds sometimes are made successfully without
preheat, especially if the carbon content of the alloy is less than 0.10 percent. In general,
preheat is preferred. Heating in the range 600-1100°F (315-593°C) is avoided because it
will result in a loss of ductility and impact strength. Welding of castings in the annealed
condition is preferred to welding of as-cast metal.
Section Size
Section size usually is considered unimportant in welding this alloy. For section
thicknesses over two inches, preheat should be above 400°F (204°C).
Cavity Dimensions
Cavity dimensions are not critical. A minimum included angle of 30° (included angles up
to 90° sometimes are used) should be maintained between the sides of the cavity, and a
root radius of 3/16 to 1/4 inch should be provided to allow full access to the root.
Welding Technique
Surfaces to be welded should be dry and cleaned to remove any residue from cavity or
weld groove preparation or other previous operations. Lack of attention to this may result
4
5
6
7
8
9
in defective welds. Either stringer or weave bead placement is used. Weaving, if any, is limited to about
the diameter of the gas nozzle. No peening is done. It is customary to remove any defects in the weld by
grinding before laying down the next bead. If a defect penetrates through the casting, or if parts to be
fabricated fit together poorly, a 3/16 inch backing plate is formed to the inside contour of the casting and
tack welded in place. The backing plate, which should be removed after welding, is generally of such a
size that it extends a minimum of 3/16 inch beyond the edge of the cavity in all directions. Tack welding
should be performed after the casting has been preheated in order to minimize the possibility of initiating a
crack at the tack weld (6 and 7).
10
Electrical Characteristics
Welding is done using DC reverse polarity. Wire diameter range is from 0.035 to 0.094
inch. Currents and voltages suggested by the manufacturer's specifications for the wire
size used are normally followed. Shielding gas is usually argon plus two percent (2%)
oxygen at a flow rate of 30 to 50 cfh. An alternate mixture of 75 percent argon plus 25
percent carbon dioxide at a flow rate of 20 cfh also is used.
Technique for Welding Machined Castings
This process is seldom used to weld machined castings; when it is, AWS ER309 or AWS
ER310 type electrode wire is used.
Post-Weld Heat Treatment
Welds usually are heated to the range 1100-1450°F (593-788°C), and then either air or
furnace cooled depending on the specification of mechanical properties for the casting. In
some cases where welds are large or located in critical areas of the casting, they are
given a full re-heat treatment of heating to 1800°F (982°C) minimum, followed by air
cooling and tempering at the specified temperature. Minor, superficial welds sometimes
are not post-heat treated when the presence of hard spots resulting from untempered
martensite in the weld deposits can be tolerated.
Non-Destructive Tests
Welds are tested for quality by one or more of the following methods of inspection:
Visual, dye penetrant, magnetic particle, radiography, pressure, or ultrasonic.
Summary
To produce welds that will satisfy the user's requirements, take the following precautions:
1.
2.
3.
Make sure that all defects have been removed to sound base metal (4), and that
surfaces to be welded are thoroughly cleaned (5 and 9).
Use the proper filler metal (2).
Use a welding technique (9) which will produce welds free of porosity,
undercutting or lack of penetration.
11
12
13
14
Gas Tungsten-Arc (GTAW)
Procedure followed by experienced producers of high alloy castings in welding
of types CA15 and CA40 alloys as reported in a survey of SFSA members
Section
1
2
3
Subject/Procedure
Base Metal
Alloy type CA15 (11.5-14 Cr, 0.15 max. C) static and centrifugal castings.
Filler Metal
AWS ER410 - Bare wire is used to weld this alloy.
Position
Whenever possible, all welding is done in the "flat" position. A ±15° angle of the groove
with the horizontal plane normally is considered flat. Successful welds can be made by
this process, however, in all positions.
Base Metal Preparation for Repair
Defects are removed before attempting any repair. Defect removal is accomplished by
grinding. Defect removal to sound base metal is assured by the use of one or more of the
following inspection processes: Visual, dye penetrant, or radiography.
Base Metal Preparation for Fabrication
Parts to be fabricated by welding are shaped to provide a groove when placed together.
The mating areas are either cast to shape and then ground, or ground or machined so
that a good fit of the welding groove can be obtained. Good practice is to machine dry
with no lubricant. Components are thoroughly cleaned before assembly. Alcohol and
acetone are solvents frequently used for cleaning.
Preheat Temperature
Heat this alloy to the range 300-600°F (149-315°C) and maintain the metal above 300°F
(149°C) during the welding operation. Welds sometimes are made successfully without
preheat, especially if the carbon content of the alloy is less than 0.10 percent. In general,
preheat is preferred. Heating in the range 600-1100°F (315-593°C) is avoided because it
will result in a loss of ductility and impact strength. Welding of castings in the annealed
condition is preferred to welding of as-cast metal.
Section Size
Section size usually is considered unimportant in welding this alloy.
Cavity Dimensions
This process is used mainly for surface welds, hence very little metal excavation is
necessary and dimensions are not critical.
Welding Technique
Surfaces to be welded should be dry and cleaned to remove any residue from cavity or
weld groove preparation or other previous operations. Lack of attention to this may result
in defective welds. Either stringer or weave bead placement is used. Weaving, if any, is
not restricted in extent. Peening may be done between successive passes on deep
welds. If parts to be fabricated fit together poorly, a 3/16 inch backing plate is formed to
the inside contour of the casting and tack welded in place. The backing plate, which
should be removed after welding, is generally of such a size that it extends a minimum of
3/16 inch beyond the edge of the cavity in all directions. Tack welding should be
4
5
6
7
8
9
performed after the casting has been preheated in order to minimize the possibility of
initiating a crack at the tack weld (6).
10
Electrical Characteristics
Welding is done using DC straight polarity. A non-consumable electrode made of
thoriated tungsten (EWTh-2) is used. A high frequency method of starting the arc is
preferred over a "scratch start" to avoid tungsten contamination of the weld. The arc
should not be struck on a carbon block. Currents and voltages suggested by the
manufacturer's specifications for the electrode size used normally are followed. Where
filler metal is used, wire sizes range from 1/16 to 3/16 inch. Either helium or argon may
be used for the inert shielding gas, but argon is preferred with a flow of 20 to 50 cfh.
Technique for Welding Machined Castings
No special technique (9) is necessary for welding machined castings; it is good practice,
however, to use small rods and low heat to avoid distortion.
Post-Weld Heat Treatment
Welds usually are heated to the range 1100-1450°F (593-788°C), and then either air or
furnace cooled depending on the specification of mechanical properties for the casting. In
some cases where welds are large or located in critical areas of the casting, they are
given a full re-heat treatment of heating to 1800°F (982°C) minimum, followed by air
cooling and then tempering at the specified temperature. Minor, superficial welds often
are not post-heat treated.
Non-Destructive Tests
Welds are tested for quality by one or more of the following methods of inspection:
Visual, dye penetrant, magnetic particle, radiography, pressure, or ultrasonic.
Summary
To produce welds that will satisfy the user's requirements, take the following precautions:
1.
2.
3.
Make sure that all defects have been removed to sound base metal (4), and that
surfaces to be welded are thoroughly cleaned (5 and 9).
Use the proper filler metal (2).
Use a welding technique (9) which will produce welds free of porosity,
undercutting or lack of penetration.
11
12
13
14
Shielded Metal-Arc (SMAW)
Procedure followed by experienced producers of high alloy castings in welding
of type CB7Cu alloy as reported in a survey of SFSA members
Section
1
Subject/Procedure
Base Metal
Alloy types CB7Cu-1 (15.5-17.0 Cr, 3.6-4.6 Ni, 2.5-3.2 Cu, 0.07 max. C) and CB7Cu-2
(14.0-15.5 Cr, 4.5-5.5 Ni, 2.5-3.2 Cu, 0.07 max. C) static and centrifugal castings.
Filler Metal
AWS E630-15
AWS E630-16
3
Lime coated electrode is preferred for DC welding. This rod
should not be used for AC.
Titania coated electrode is used for AC welding and may be used
for DC.
2
Position
All welding is done in the "flat" position. A ±15° angle of the groove with the horizontal
plane normally is considered flat.
Base Metal Preparation for Repair
Defects are removed before attempting any repair. Defect removal is accomplished by
arc-air, chipping, gouging, grinding, or machining, or by some combination of these
operations. Defect removal to sound base metal is assured by the use of one or more of
the following inspection processes: Visual, dye penetrant, or radiography. Where dye
penetrant or radiographic inspection of a prepared cavity discloses shrinkage of a severity
not in excess of that specified for the casting as a whole, acceptable practice is to weld
such areas without further preparation (3.2.1).
Base Metal Preparation for Fabrication
Parts to be fabricated by welding are shaped to provide a groove when placed together.
The mating areas are either cast to shape and then ground, or ground or machined so
that a good fit of the welding groove can be obtained. Good practice is to machine dry
with no lubricant. Components are thoroughly cleaned before assembly. Alcohol and
acetone are solvents frequently used for cleaning.
Preheat Temperature
Normally this alloy is not preheated; however, if the section size is over 3/4 inch in
thickness, and the extent of the weld substantial, the alloy may be preheated to 500°F
(260°C).
Section Size
Section size usually is considered unimportant in welding this alloy. Thick sections may
require preheat (6) for satisfactory welds.
Cavity Dimensions
Cavity dimensions are not critical. A minimum included angle of 30° (included angles up
to 90° sometimes are used) should be maintained between the sides of the cavity, and a
root radius of 3/16 to 1/4 inch should be provided to allow full access to the root.
Welding Technique
Surfaces to be welded should be dry and cleaned to remove any residue from cavity or
weld groove preparation or other previous operations. Lack of attention to this may result
4
5
6
7
8
9
in defective welds. Either stringer or weave bead placement is used. Weaving, if any, is
limited to two and one-half times the electrode diameter. Fully hardened castings are
frequently preheated (6) and welded with low heat and small rods. No peening is done.
All slag is removed with a stainless steel wire brush or slagging hammer, or needle gun
using stainless steel needles. If a defect penetrates through the casting, or if parts to be
fabricated fit together poorly, a 3/16 inch backing plate is formed to the inside contour of
the casting and tack welded in place. The backing plate, which should be removed after
welding, is generally of such a size that it extends a minimum of 3/16 inch beyond the
edge of the cavity in all directions. Tack welding should be performed after the casting
has been preheated in order to minimize the possibility of initiating a crack at the tack
weld (6 and 7).
10
Electrical Characteristics
Welding normally is done using DC reverse polarity. Successful welds can be made,
however, using AC. Electrode sizes from 3/32 to 1/4 inch may be used with the
amperage and voltage suggested by the electrode manufacturer's specifications for the
particular size rod. Due to the high electrical resistance of stainless steel, the burn-off
rate of the electrode is higher than for carbon steel. Arc length should be maintained as
short as possible. A short arc length is very important when starting a weld pass since a
long arc can sometimes be caused by initial hand recoil and may result in weld spatter or
porosity.
Technique for Welding Machined Castings
No special technique (9) is necessary for welding machined castings; it is good practice,
however, to use small rods and low heat to avoid distortion. If the welded area will be
subject to corrosion, it is desirable to quench the weld zone with a wet cloth between each
pass. For small welds on heavy sections, this may not be necessary since the heavy
mass will tend to cool the weld zone rapidly.
Post-Weld Heat Treatment
Both annealed and aged type CB7Cu castings can be restored to specified hardness by
low temperature postweld hardening treatment in the range 900-1100°F (482-593°C). But
to restore hardenability properties to multiple-pass welds on heavy sections, they are
heated to the range 1850-1950°F (1010-1066°C), held until uniformly at temperature,
rapidly cooled by quenching in water, oil or air, and followed by the desired aging
treatment. Single-pass welds usually do not require postweld solution heat treatment.
Non-Destructive Tests
Welds are tested for quality by one or more of the following methods of inspection:
Visual, dye penetrant, radiography, or pressure.
Summary
To produce welds that will satisfy the user's requirements, take the following precautions:
1.
2.
3.
Make sure that all defects have been removed to sound base metal (4), and that
surfaces to be welded are thoroughly cleaned (5 and 9).
Use the proper filler metal (2).
Use a welding technique (9) which will produce welds free of porosity,
undercutting or lack of penetration.
11
12
13
14

metallurgical defect in ductile iron

2009-09-01T23:50:00.000-07:00

در اين قسمت مقاله اي در رابطه با عيوب متداول در قطعات چدن خاكستري را مورد بحث و بررسي قرار خواهيم داد.در صورت نياز به كل مقاله و يا مقالات مرتبط در اين رابطه تماس بگيريد.

Metallurgical Defects in Ductile Cast Iron
Causes and Cures
C.M.Ecob
Customer Services Manager, Elkem AS, Foundry Products Division
Abstract
The objective of this paper is to provide an overview of some of the most common
metallurgical defects found in the production of ductile cast iron today. The examples shown
have all been determined during the examination of samples in Elkem’s Research facility in
Norway.
Whilst many foundries recognise the defects, an appreciation of the possible causes, and
therefore cures, is not always apparent. The causes and cures for the different problems are
examined in the paper. Emphasis is made on shrinkage problems, probably the most common
problem seen by Elkem’s team of service engineers around the world.
Introduction
Metallurgical defects in ductile iron can be very costly to the foundry, not only because the
part has to be remade or rectified, but due to the unfortunate fact that many defects are not
revealed until after the expensive machining stage. Care in the selection of raw materials,
good process control in the melting stage and proper metal handling procedures will go a long
way to the prevention of defects.
Further, a routine for logging and recording of defect occurrences will reveal which are the
major problem areas, allowing for a systematic elimination of the defects.
This paper will examine the most common defects, starting with shrinkage. Deterioration of
affordable steel scrap qualities, use of incorrect inoculants and nodularisers plus the pressures
to get castings out of the door as fast as possible has led to an increase in the incidences of
shrink/porosity related cases seen by Elkem’s team of technical service engineers. Indeed, the
ductile iron foundry, which truthfully claims not to have shrinkage concerns is the exception
to the rule.
Other common defects may be divided into two basic categories:
-Those related to nodule shape and size, such as compacted graphite structures,
exploded and chunky graphite, graphite floatation, spiky graphite and nodule alignment.
-Those related to inclusions/abnormalities within the matrix, such as flake graphite
surfaces, slag inclusions, carbides and gas.
These problem areas are described to aid recognition of the defect and causes are discussed
together with possible cures.

freeze casting prototyping(invest ment casting)

2009-08-21T23:26:00.000-07:00

در اين پيام مقاله اي در رابطه با ريخته گري دقيق با مدل يخي را كه از اخرين دستاوردهاي علمي جامعه ريخته گران به شمار مي رود را برايتان معرفي مي نمايم.در صورت نياز بع=ه اطلاعات بيشتر تماس بگيريد.

تحولي نوين در صنعت ريخته گري : ريخته گري دقيق با مدل يخي

واحد نجاتي مازگر
شركت ريخته گري تراكتورسازي ايران - تبريز

اخيرا محققان از مدل يخي جهت ساخت قالب سراميكي براي ريخته گري دقيق استفاده به عمل آورده اند.در اين روش بدليل مشخصه هاي ويژه يخ ،مخلوط قالب سراميكي بايستي به گونه اي باشد كه در دماي زير صفر درجه سانتي گراد به خود گيري لازم برسد و اين متفاوت از قالبهاي پوسته اي مورد استفاده در ريخته گري دقيق با مدل هاي مومي مي باشد كه در دماي محيط و بالاتر به خودگيري لازم مي رسند.مهم ترين پارامتري كه در ريخته گري با مدل يخي بايستي كنترل گردد تافنس شكست قالب سراميكي مي باشد كه در واقع مقاومت و استحكام آن را در مراحل بعدي ذوبريزي نشان مي دهد.مخلوط قالبگيري در اين روش عموما حاوي سيليكاي گداخته ،پودر سيليكات آلومنيوم، رزين و كاتاليزور مي باشد كه در مراحل بعدي بيشتر توضيح داده خواهند شد.در اين مقاله با تغيير دادن نسبت وزني سيليكاي گداخته فيبردار به پودر سيليكات الومنيوم و نيز مقادير رزين و كاتاليست ،مقدار تافنس شكست اندازه گيري و سطع مقطع شكست نمونه ها براي بررسي مكانيزم شكست مورد بحث و بررسي قرار گرفته است.نتايج نشان مي دهد افزايش سيليكاي گداخته فيبردار باعث افزايش تافنس شكست نمونه و مقاومت در برابر ترك ناشي از ذوبريزي مي گردد.در اين ميان ميزان رزين مصرفي و كاتاليست نيز روي نسبت بهينه مورد اشاره تاثير مستقيم مي گذارد.

خلاصه:

موم معمول ترين ماده مورد استفاده در ساخت مدل براي ريخته گري دقيق مي باشد.تنش هاي اعمالي از طرف موم در حين ذوب براي تخليه شدن باعث بروز مشكلات عمده اي شامل بروز ترك در پوسته مي گردد[1].مطابق آنچه كه ريچارد و ساندرلند [2] نشان داده اند،موم تا حدي مشابه پليمرهاي خطي بلوري مي باشد كه در اثر ذوب شدن يك افزايش ناگهاني در حجم آن رخ مي دهد.اين افزايش حجم عامل اصلي بروز ترك در قالب پوسته اي در مرحله پخت به شمار مي رود.در شكل 1 مقايسه اي از پوسته توليد شده با مدلهاي يخي و مومي نشان داده شده است.
شكل 1) مقايسه اي از بروز ترك در پوسته ساخته شده از مدل يخي نسبت به مدل مومي

همانطور كه در شكل 1 ملاحظه مي گردد تمايل به بروز ترك در پوسته توليدي از مدل مومي بالا مي باشد و اين در حالي است كه ،ميزان ترك در پوسته حاصل از مدل يخي در حد صفر مي باشد.
فرايند ريخته گري با مدل يخي (FCP)[1] يكي ديگر از روش هاي ريخته گري دقيق به شمار مي رود كه توسط Yodice [3] اختراع گرديده است..در اين فرايند كه اساس آن مشابه ريخته گري دقيق با مدل مومي مي باشد از يخ در ساخت مدل براي ريخته گري قطعات فلزي استفاده به عمل مي آيد.Yodice و ساير همكارانش امكان پذيري اين روش را كاملا به اثبات رساندند[6-4].مزيت عمده روش ريخته گري با مدل يخي براي جايگزين كردن آن با روش قديمي مدل مومي مي تواند مقرون به صرفه بودن ،سهولت فرايند و كيفيت سطحي بسيار بالا در توليد قطعات نزديك به شكل نهائي مي باشد.براي توليد مدل هاي سه بعدي و پيچيده يخي عموما از روش نمونه سازي با انجماد سريع استفاده مي كنند.در اين روش مدل يخي را مستقيما از روي مدل طراحي شده كامپيوتري و با انجماد لايه به لايه قطره هاي آب مي سازند[10-7].
دراين روش با توجه به ذوب شدن يخ در دماي محيط فرايند قالبگيري مي بايستي در دماي زير صفر صورت گيرد.در نتيجه مواد قالبگيري در اين روش بايستي مشخصه هاي ويژه اي داشته باشند و اين نياز به مطالعات وسيعي براي يافتن موادي كه در دماي زير صفر به خودگيري لازم برسند ،دارد.امروزه اگرچه اين روش به خوبي قابل اجرا مي باشد ولي تعداد منابع علمي و تدوين شده در اين رابطه بسيار محدود مي باشد و هنوز جوانب ناشناخته و پنهان زيادي در اين رابطه وجود دارد.بررسي ها نشان مي دهد موفقيت در انجام روش FCP شديدا وابسته به تافنس(چقرمگي) قالب ساخته شده مي باشد.لذا مطالعات روي اين پارامتر همراه با بررسي سطوح شكست اساس اين تحقيق را تشكيل مي دهد..پارامترهاي اصلي كه در اين تحقيق مورد بررسي قرار گرفته است شامل نسبت وزني سيليكاي گداخته[2](FS) حاوي فيبر به ماده نسوز سيليكات آلومنيوم و مقدار حجمي رزين و كاتاليست مورد استفاده در تهيه مخلوط قالبگيري مي باشد.همچنين براي تبيين مكانيزم شكست در قالبهاي پوسته اي سطح شكست و ريزساختار نمونه هاي آزمايشي با ميكروسكوپ الكتروني تحت بررسي قرار گرفته شد.مدل انتخابي در اين تحقيق پيچ M8 مي باشد تا از نقاط تيز و پيچيدگي لازم برخوردار باشد.
انتخاب مواد مناسب براي قالبگيري در دماي زير صفر:
موادي كه براي ساخت قالب در ريخته گري دقيق با مدل يخي مورد استفاده قرار مي گيرد بايستي مشخصه هاي ويژه و متفاوت با آنچه در مدل مومي مورد نياز بود را داشته باشد[15-11].براي مثال چسب مورد استفاده در روش FCP بايستي عاري ازآب باشد تا در دماي زير صفرمنجمد نگردد.حال در اين قسمت مخلوط قالبگيري براي توليد پوسته سراميكي را توضيح خواهيم داد.
مخلوط سراميكي:
همانند ديگر روش هاي ريخته گري دقيق دوغاب مورد استفاده براي پوشش دهي مدل يخي مخلوطي از چسب و مواد نسوز سراميكي مي باشد.لذا اين مواد نقش حياتي را در ساخت پوسته ايفا مي كنند.دامنه تغييرات مواد نسوز مورد استفاده بسيار وسيع مي باشد و معمولا تركيبي از ماسه سيليسي ،سيليكات آلومنيوم،آلومين،سيليكاي گداخته(FS) و سيليكات زيركونيوم مي باشد.انتخاب مواد نسوز مناسب در حصول قطعه اي با كيفيت سطحي ،دقت ابعادي و خواص بالا امري ضروري مي باشد.
در انتخاب مواد قالبگيري موارد زير بايستي مد نظر قرار گيرد:درجه ديرگدازي،ضريب انبساط حرارتي ،تركيب شيميائي،مسايل هزينه اي،اندازه و توزيع دانه بندي مواد نسوز.
همانند ماسه ماهيچه و قالبگيري دانه بندي ذرات مواد نسوز و نيز نسبت آنها در تركيب مواد قالبگيري روي كيفيت نهائي پوسته و قطعه حاصله مستقيما اثرگذار مي باشد.در اين تحقيق از سيليكات آلومنيوم به عنوان ماده نسوزقالبگيري استفاده بعمل آمد.ولي همانطوركه درشكلa2نشان داده شده است استحكام حاصله در پوسته پائين بوده و باعث ايجاد ترك در ان شده است.براي بهبود آن از وجود سيليكاي گداخته فيبردار در تركيب مواد قالبگيري استفاده به عمل آمد و همانطور كه در شكل b 2 نيز نشان داده شده است باعث بهبود استحكام در قالب پوسته اي شده است.
شكل 2) پوسته سراميكي(a) بدون / (b) با حضور سيليكاي گداخته فيبردار

چسب:
جزء اصلي ديگر در مخلوط قالبگيري چسب مي باشد كه در اثر تركيب با مواد نسوز تشكيل دوغاب را مي دهد.براي ريخته گري دقيق در دماي پائين حلال مورد استفاده در چسب نبايستي منجمد گردد و در دماي زير صفر نيز از سياليت لازم برخوردار بوده باشد .در كنار اين مشخصه ها لازم است قالب حاصل از آن استحكام بالائي داشته و باعث حصول قطعه اي با كيفيت سطحي و دقت ابعادي عالي گردد.بر اساس معيار هاي اشاره شده در اين تحقيق سه نوع چسب مورد آزمايش قرار گرفت:آب شيشه،سيليكاي ژله اي و سيليكات اتيل.قالب ساخته شده با چسب سيليكات سديم (آب شيشه)داراي استحكام گرم پائيني بود و علاوه بر آن دقت ابعادي و كيفيت سطحي قطعات توليدي نيز پائين مي بود.به همين دليل در آزمايشات بعدي از آب شيشه استفاده نگرديد.چسب دوم سيليكاي ژله اي بود كه از ترقيق كردن آب شيشه به ميزان 10-9 برابر با افزودن آب بدست مي آيد.لذا با توجه به اينكه 50-40 درصد وزني آن را آب تشكيل مي دهد در دماي زير صفر سريعا منجمد گرديده و لذا امكان استفاده از سيليكاي ژله اي نيز مقدور نگرديد.در مرحله بعدي آزمايشات كه از سيليكات اتيل استفاده به عمل آمد تمامي خواص مورد نياز حاصل گرديد.
كاتاليست:
جهت كوتاه كردن مدت زمان ژله اي شدن و خودگيري دوغاب (مخلوط چسب با مواد نسوز سراميكي) جهت ساخت قالب براي ريخته گري دقيق در دماي پائين نياز به استفاده از يك كاتاليزور مي باشد.به عبارتي ديگر بدون استفاده از كاتاليزور ممكن است زمان ژله اي شدن روزها به طول انجامد.مدت زمان ژله اي شدن به عواملي چون دما ،تركيب چسب مصرفي و PH دوغاب بستگي دارد.اثر PH در اين رابطه برجسته تر از دو عامل ديگر مي باشد.زماني كه عدد PH برابر با 2 باشد پايداري دوغاب افزايش يافته و در نتيجه مدت زمان ژله اي شدن افزايش مي يابد.زماني كه6-5 PH= باشد دوغاب بسيار ناپايدارشده و درنتيجه زمان ژله اي شدن كوتاهتر مي گردد.زماني كه 1 PH< باشد نيز شرايط مشابهي پيش مي آيد.لذا كاتاليزوز مورد استفاده بايستي PH دوغاب را از دامنه پايدار به ناپايدار تغيير دهد.لذا با توجه به استدلال بالا كاتاليزور مورد استفاده يا بايستي خاصيت اسيدي داشته باشد كه عدد PH را به زير يك برساند و يا اينكه خاصيت قليايي داشته باشد تا عدد PH را به دامنه 7-6 برساند.در اين ميان از جمله مهم ترين كاتاليست هاي اسيدي مي توان به اسيد سولفوريك،كلريك و فسفريك اشاره نمود و از جمله مهم ترين كاتاليست هاي بازي نيز مي توان به Ca(OH)₂ Na(OH), MgO, Mg(OH)₂,Na₂CO₃ اشاره نمود.با توجه به نتايج آزمايشات انجام پذيرفته در مرجع [6] استفاده از كاتاليزورهاي آلي باعث تشكيل دوغاب و در نتيجه قالب پوسته اي با كيفيت و خواص بالا و قطعه با كيفيت سطحي عالي مي گردد.در مقابل كاتاليزور هاي اسيدي در مرحله پخت به سطح داخلي پوسته رفته و باعث افت شديد خواص مكانيكي در دماي بالا مي گردد.همچنين بقاياي كاتاليستهاي غير آلي باعث خروج فازهاي پيوندي از حالت شيشه اي و در نتيجه بروز ترك در قالب مي گردد.اين در حالي است كه كاتاليستهاي آلي در مرحله پخت كاملا سوخته و از سيستم خارج شده و لذا خواص گرم پوسته را تحت تاثير قرار نمي دهد.
مخلوط دوغاب:
اين مخلوط شامل پودر سراميكي ،چسب و كاتاليست مي باشد.درصد وزني هر كدام از اين اجزا روي عملكرد دوغاب ،كيفيت و خواص قالب پوسته اي و در نهايت كيفيت قطعه نقش مهمي را ايفا مي نمايد.بر اين اساس مواد زير براي انجام ازمايشات انتخاب گرديدند:
مخلوطي از سيليكات آلومنيوم و سيليكاي گراخته فيبردار به عنوان پودر سراميكي و سيليكات اتيل هيدروليز شده به عنوان چسب و مخلوطي از اتانول با تري اتانول آمين (با نسبت 1.5:1) به عنوان كاتاليست .
دانه بندي و توزيع ذرات سيليكاي گداخته و سيليكات آلومنيوم در جداول 2و1 آورده شده است.

جدول 1) دانه بندي و توزيع ذرات سيليكاي گداخته مورد استفاده

جدول 2) دانه بندي و توزيع ذرات سيليكات آلومنيوم مورد استفاده
جهت اطمينان از تكرار پذيري نتايج آزمايشات اقدامات احتياطي زير انجام پذيرفت:
· نگهداري چسب در داخل شيشه هاي ايزوله و قرار دهي آن در يخچال جهت رسيدن به دماي محيط كاري (℃ 15- )
· نگهداري پودرهاي سراميكي در ظروف شيشه اي ايزوله و قرار دهي آن در يخچال جهت رسيدن به دماي محيط كاري
حرارت دادن آنها به مدت 2 ساعت در 240 درجه سانتي گراد جهت كنترل رطوبت قبل از استفاده
· اختلاط تري اتانول امين و اتانول با نسبت هاي ذكر شده و نگهداري آن در شيشه هاي ايزوله و قرار دهي آن در يخچال جهت رسيدن به دماي محيط كاري
· اندازه گيري دقيق وزني هر كدام از مواد در هر آزمايش

روش انجام آزمايش :

متغيرهاي فرايند:
متغير هاي اين فرايند بدين شرح انتخاب گرديد:مقدار كاتاليست،مقدار چسب و نسبت سيليكاي گداخته به سيليكات آلومنيوم.همانطور كه در جدول 3 نيز نشان داده شده است براي هر كدام از اين متغيرها نيز سه داده در نظر گرفته شد.

جدول 3) مقادير انتخاب شده براي متغيرهاي آزمايش
لذا همانطور كه در جدول 4 نيز نشان داده خواهد شد تحقيق حاضر شامل 9 آزمايش خواهد بود.لازم به ذكر است زمان اختلاط با توجه به تجربيات قبلي نويسنده روي 4 دقيقه ثابت و دماي كاري در تمامي آزمايشات ℃15- در نظر گرفته شد.

آماده سازي نمونه ها:
نمونه هاي آزمايش مورد استفاده در اين آزمايش به روش ريخته گري دوغاب تهيه گرديد.بدين شكل كه دوغاب آماده شده به شرح زير در داخل قالب يخي ريخته شده و سپس با ذوب كردن يخ نمونه هاي بدست آمدند.
1- پودر هاي از پيش خشك شده به محض خروج از فريزر وزن شده و در داخل مخزن ميكسر تحت خلا ريخته مي شود.
2- چسب توسط ترازوي از پيش سرد شده توزين شده و در داخل مخزن ميكسر تحت خلا ريخته مي شود.
3- كاتاليست از پيش مخلوط شده به داخل مخزن ميكسر تحت خلا ريخته مي شود.
4- سرپوش مخزن ميكسر گذاشته مي شود و عمليات اختلاط تحت شرايط خلا صورت مي پذيرد.
5- بلافاصله پس از اتمام اختلاط دوغاب حاصله در داخل قالب يخي ريخته مي شود.
لازم به ذكر است پس از اتمام اختلاط و قبل از بارريزي دماي مخلوط به ℃4- رسيده بود.
در شكل 3 ابعاد و شكل هندسي نمونه هاي آزمايشي ريخته شده در قالب يخي بر اساس استاندارد ASTM C1161-02c نشان داده شد ه است.

شكل 3) ابعاد و شكل هندسي نمونه هاي آزمايشي مطابق استاندارد ASTM C1161-02c
پس از پخت و درآوردن نمونه ها ابعاد آنها توسط دستگاههاي دقيق اندازه گيري شد و سپس جهت تست خمشي با دستگاه داراي چهار تكيه گاه آماده گرديد.در شكل 4 نماي شماتيك دستگاه تست خمش همراه با توزيع تنش در آن نشان داده شده است.
شكل 4) نماي شماتيك دستگاه تست خمش و نمودار توزيع تنش در آن

در رابطه با دستگاه تست خمش:
تست خمش با مشخصات ذكر شده براي دستگاه آن سالهاست كه روش قابل اعتمادي براي اندازه گيري استحكام خمشي به شمار مي رود.در اين دستگاه مبناي اندازه گري بر اساس رابطه زير مي باشد:
كه در اين رابطه S استحكام خمشي ، P نيروي اعمالي براي شكست،L فاصله بين دو تكيه گاه بيروني ،b عرض و d ضخامت نمونه آزمايشي مي باشد كه در شكل 4و3 نشان داده شده است.براي محاسبه تافنس شكست سطح زير منحني توزيع تنش محاسبه مي گردد كه نمونه اي از آن در شكل 4 نشان داده شده است.لازم به ذكر است نرخ بارگذاري توسط دستگاه 10psi/min مي بود و براي هر كدام از آزمايشات طراحي شده در جدول 4 به تعداد 20 عدد نمونه آزمايش گرديد و در مجموع 180 عدد نمونه ريخته شد.
جدول 4) شرايط آزمايشات طراحي شده براي اين تحقيق
فرآيند توليد قطعه در FCP :
فرايند توليد قطعه در روش ريخته گري دقيق با مدل يخي تا حد زيادي مشابه همين روش با مدل مومي بوده و شامل مراحل زير مي باشد:
· ساخت مدل يخي با استفاده از قالب سيليكني و يا روش نمونه سازي سريع انجمادي
· توليد سيستم راهگاهي از يخ به روش مشابه و مونتاژ آنها به مدل مربوطه
· پوشش دهي مجموعه مونتاژ شده با يك ماده جداكننده [14-11]
· دوغاب دهي مجموعه حاصله با دوغاب سراميكي از قبل آماده شده
· انتقال مجموعه دوغاب دهي شده پس از ژله اي شدن در دماي محيط جهت ذوب شدن يخ
· حرارت دهي در دماي ℃900 به مدت زمان معين براي شكل گيري قالب پوسته اي
· ريختن مذاب با شرايط معين در داخل پوسته سراميكي
· در آوردن قطعه ريختگي و عمليات تميزكاري بعدي
بحث و نتيجه گري:
نتايج استحكام خمشي حاصل از تست شكست نمونه هاي سراميكي نشانگر وجود پراكندگي قابل ملاحظه داده ها مي باشد.جهت براورد ميزان قابليت اطمينان به نتايج حاصله نياز به يكسري ابزارهاي آماري جهت تعيين و مقايسه تغييرپذيري مي باشد.براي اين منظور مدول Weibull در اين تحقيق مورد استفاده قرار گرفت كه مجال پرداختي به اين موضوع در اين مقاله ميسر نمي باشد.لازم به ذكر است در جدول4 اين مشخصه براي هر كدام از آزمايشات تعيين گرديده است.
نتايج بدست امده از اين تحقيق نشانگر اين موضوع مي باشد كه با افزايش مقدار سيليكاي گداخته فيبردار تافنس شكست پوسته سراميكي زماني افزايش مي يابد كه مقدار سيليكاي گداخته افزوده شده از حد معيني بيشتر باشد.همچنين با كاهش ميزان چسب مصرفي تافنس شكست افزايش مي يابد.بررسي هاي انجام پذيرفته با ميكروسكوپ الكتروني روي سطوح شكست نشانگر اين موضوع مي باشد كه افزودن سيليكاي گداخته فيبردار باعث ايجاد يكپارچگي ساختاري و ممانعت از بروز ترك در آن مي گردد.در شكل 5 اين موضوع به وضوح نشان داده شده است.

a b
شكل 6) مقايسه اي از ريزساختار يكپارچه a (داراي فيبر) و تركدار b (بدون فيبر) در تركيب سيليكاي گداخته
دقت ابعادي و كيفيت سطحي بالاي نمونه هاي ريخته شده نشانگر اين موضوع مي باشد كه ريخته گري دقيق با مدل يخي براي توليد قطعات با پيچيدگي هندسي بالا بسيار مناسب مي باشد.در شكل 6 نمونه اي از پيچ هاي M8 ريخته شده در اين تحقيق نشان داده شده است.

شكل 6) نمونه اي از قطعات آزمايشي ريخته شده

فهرست منابع و مراجع :
مقاله حاضر برگردان مقاله زير مي باشد كه همراه با آن منابع و مراجع ديگري نيز براي مطالعه بيشتر آورده شده است.
- Fracture toughness of ceramic moulds for investment casting with ice patterns
Q. B. Liu*, M. C. Leu and V. L. Richards

1. P. R. Beeley and R. F. Smart: ‘Investment casting’, 1st edn; 1995,Cambridge, The University Press.
2. V. L. Richards and B. V. Sunderland: Proc. 49th Annual Technical Meeting, rlando, FL, USA, October 2001, Investment Casting Institute, 2:1–12.
3. A. Yodice: US patent 5,072,770, 1991.
4. A. Yodice: INCAST: Int. Mag. Invest. Cast. Inst., 1998, 11, 19–21.
5. D. M. Peters: Foundry Manage. Technol., 1995, 123, (90–91),96.
6. X. Wu: ‘Study on ceramic mold investment casting based on ice patterns made by rapid prototyping method’, Bachelor’s dissertation,Tsinghua University, Beijing, hina, 1997.
7. M. C. Leu, Q. Liu and F. D. Bryant: Ann. CIRP, 2003, 52, 185–188.
8. Q. Liu, G. Sui and M. C. Leu: J. Comput. Ind., 2002, 48, 181–197.
9. Q. Liu and M. C. Leu: Proc. International Mechanical Engineers Congress and Exposition, Anaheim, CA, USA, November 2004,American Society of Mechanical Engineering, Paper IMECE2004-59580.
10. W. Zhang, M. C. Leu, Z. Ji and Y. Yan: US patent 6,253,116, 2001.
11. Q. Liu, M. C. Leu, V. L. Richards and S. M. Schmitt: Int. J. Adv.Manuf. Technol., 2004, 24, 485–495.
12. Q. Liu and M. C. Leu: Proc. Conf. on ‘Materials and process for medical devices’, Anaheim, CA, USA, September 2003, ASM International, 438–443.
13. M. C. Leu and Q. Liu: Proc. NSF Design, Service and Manufacturing Grantees and Research Conference, Birmingham,AL, USA, January 2003, The University of Alabama.
14. Q. Liu and M. C. Leu: Proc. Symp. on ‘Solid freeform fabrication’,Austin, TX, USA, August 2002, University of Texas at Austin,563–574.
15. Q. Liu, M. C. Leu, V. Richards and H. Jose: Proc. 15th Symp. On ‘Solid freeform fabrication’, Austin, TX, USA, August 2004,University of Texas at Austin, 602–611.
16. ‘Standard test method for flexural strength of advanced ceramics at
ambient temperature’, C 1161-02c, ASTM, Philadelphia, PA, USA,2002.
17. F. I. Baratta: ‘Requirements for flexure testing of brittle materials’,AMMRC TR 82-20, Army Materials and Mechanics Research Center, Watertown, MA, USA, 1982.18. V. L. Richards and G. Connin: Proc. 49th Annual Technical Meeting, Orlando, FL, USA, October 2001, Investment Casting Institute, Paper 13.
19. D. R. Askeland: ‘The science and engineering of materials’, 3rd edn; 1989, Boston, MA, PWS Publishing Company.
20. K. Kendall, N. McN Alford, S. R. Tan and J. D. Birchall: J. Mater. Res., 1986, 1, 120–123.
21. V. L. Richards and S. Mascreen: Proc. 50th Technical Conference and Expo, Chicago, IL, USA, September–October 2002, Investment Casting Institute.
22. G. E. Dieter: ‘Mechanical metallurgy’, 3rd edn; 1986, New York, McGraw-Hill Book Company.

[1] - Freeze Cast Process
[2] - Fused Silica

مقلات جديد در رابطه با جوانه زني چدنها

2009-08-15T22:02:00.000-07:00

در صورت نياز به هر كدام از مقالات تماس بگيريد

- Inoculation of Cast Iron
- Inoculant Alloy Composition
- Inoculation Practices
- Inoculation Mechanisms
- Fading of Inoculation

معرفي مقالات جديد

2009-08-11T21:53:00.000-07:00

در صورت نياز به هر كدام از مقالات زير تماس بگيريد:

-Effect of graphite inclusion on mechanical properties of ADI

-Influence of austempering variable on dimensional chanche of Cu-Ni D.I

-Effect of semi austempering on the fatique properties of DI

-Ductile iron pipe system-

carbon in green sand molding

2009-08-08T20:59:00.000-07:00

The Significance of Total Carbon in Green Sand Systems
Controlling the quality of raw materials and additives is critical to producing high-quality ferrous castings — as is testing and managing these factors, and the production processes to enhance productivity.
Basic understanding of raw materials is equally important as good quality control measures and consistent operating practice.
Controlling raw materials is critical to the success of iron casting production from green sand systems. The base silica sand is often overlooked, with the main focus on bentonite additions. Carbonaceous additives can be considered a “necessary evil” to ensure a good surface finish and reduction in sand-related surface defects. Other additives are used when systems get out of balance, and these will further complicate the complexity of green sand systems. For castings requiring cores this becomes a bigger issue, as many differing resin systems are employed for core production, and these must be taken into consideration when controlling both the carbonaceous levels and the overall grading of the sand system. The twin effects on additional carbon and loss-on-ignition, and overall sand grading, need careful understanding and control.
Control of raw materials Much has been documented about additives to green sand systems. What follows here is a discussion of the need for a complete understanding of the sand system, so that operators can make judgments based on fact. Of the many green sand systems I have observed worldwide, most share many common features that give cause for concern. This is not because the operators do not have the means to control the system, but because their decisions are based on a lack of validated data. Some of my concerns are:• Poor sampling; • Frequency and timing of samples; • Lack of calibrated testing; • Focus on the wrong control areas; • Lack of understanding of primary and secondary sand tests; • Little focus on silica sand or sand grading; • Bentonite emphasizes detriment to carbonaceous additive;• Poor in-coming test procedures;• Over-reliance on suppliers. All of these are important, but equally good castings can be produced from systems with little or no control. By control, I mean a basic understanding of all the raw materials used together with in-house testing and/or approved certification, coupled with good, consistent casting practices. This is measured by general scrap rates and the costs associated with knockout and shot blasting, costs so often not taken into consideration when selecting the raw materials to be used in the system.
Mass balance checklist Return-sand storage silo/bunker capacity Return-sand temperature/moisture (at various points) Weight of sand on line/in boxes etc Core weight input (if applicable) Fines extraction/Sand losses Additive control/weigh calibration/stock against usage Sand carryover at knockout Shot blasting times/consumption Scrap levels/sand related defects Sand to metal ratio/casting weight data Volatiles data at mill and in return sand Loss-on-ignition data at mill and in return sand
A great deal of information can be gathered from doing a complete mass balance on a green sand system. This exercise picks up so much useful data that it should be a regular exercise in all foundries, especially if the casting weights or size of castings alters over time. Consider the simple question “What is the total sand weight in the green sand system?” Normally, this is met either with a wild estimate or a complete blank response.
From the data gathered a much clearer picture may emerge. Most important, we know that control systems can actually monitor the burn-out rates of bentonite and carbonaceous additives. Most observers agree that additions at the muller require time to be effective, and by understanding how the sand reacts to varying sand-to-metal ratios, proactive steps can be taken. Various foundries have used predictive software or even a simple traffic light system to monitor heat demand, with good results. It is not the actual system that matters, but, importantly, understanding what is happening and being one step ahead of it.
Those foundries producing castings without cores are obviously in a much better position than those with cored products. Core systems simply have to accept coarser sand as a dilution, but with careful selection this does not have to be a problem. More important is knowing the AFS (Average Fineness Number) and AGS (Average Grain Size), and determining what produces the best casting surface at the most economical cost. The AFS clay-grade washing and sieving of the washed sand needs regular review, and these tests, along with optimum additive rates, are the keys to success in a green sand system.
Bond suppliers will claim many advantages for their products, but many differing bentonites are used with success throughout the iron-casting world. That said, consistency is the main aspect to control, and that is another matter entirely. Few foundries have the specialized equipment needed for incoming control, so testing should be limited to meaningful values and working with your supplier will establish a working specification.
A smooth surface, free from defects is the standard for excellent casting surface finish.
Many coal products are available, with many claims. Coal, either as high-quality bituminous or lower-grade material, is the main additive, and grading selection should also be tailored to the castings being produced. Coal is by far the cheapest and safest carbonaceous additive to green sand systems. Its unique, all-around properties make it easy to use; its use is so underrated that in most cases it is under-used, and most systems do not benefit fully from its various beneficial properties.
Most other carbon additives are highly volatile products normally added to poorer quality coals to enhance the level of volatile content. These can cause additional problems in high concentrations due to hydrogen/nitrogen gases, and are quickly eliminated from a system without contributing to the coke build-up (total carbon). This can be an expensive approach to adding carbonaceous additives, and too often the operator has no idea what products he/ she is using, as the blends are often “trade secrets.” To understand a system and its interaction with additives you simply have to know what you are adding, and why. The classic system of using one, single-source bentonite and coal is by far the simplest and most cost-effective route to producing quality castings.
Mixer controls, technical support The need for control in this area is obvious as changes here are reflected throughout the system. Regular mixer maintenance coupled with calibration of additives and usage checks against actual stock purchases often may spot problems before they give concern. This should be part of the mass balance procedure that should be a routine feature.
Getting the primary and secondary testing balance right is difficult for foundries. Too much, and the testers see no positive advantages other than routine test control. Too little, and there is no starting point for investigation and correction once things go wrong.
Concentrate on the primary tests to ensure consistency. Especially during periods of good casting performance, take time to monitor what makes the system work. Look to the secondary tests to reinforce the casting quality, and test only if you are prepared to react on the results. Meaningless testing is a waste of resources. Ensure proper reporting and graphing of data. Ultimately, all sand testing shows a “trend.” Even if the testing methods are not perfect, if conducted on a consistent basis they will show the trend line to which you are able to react.
Bentonite testing has been well documented. One area for control is a foundry’s raw-material acceptance testing. This can be established in conjunction with the supplier, and simple tests such as swelling ability, sieve grading, and specific gravity are useful in association with active clay calibrations.
One of the most difficult tests from which to get consistent results is the volatiles test, and yet many foundries use this as the guide for carbonaceous additions irrespective of the casting quality. With 3 or 4 decimal place measurement and a weight loss at a specified temperate (910ºC) and time (7 minutes) on a small sample weight (1 or 2 grams), it can be a difficult test, especially between different laboratories. This volatiles test is normally used in conjunction with the loss-onignition test and investigations into an alternative method to supplement these tests will ensure focus on this critical area for consistent casting performance.
Total carbon testing For example, Lincoln Castings Ltd. began to use this control method on its LM1 and LM2 molding lines, starting around 2001. It was introduced as a routine test for their straight coal/ clay system. Over the following year it became obvious that the interaction of total carbon, sulfur, volatiles, and loss-on-ignition was critical to the quality of the castings.
The radar graph shown above was used to monitor the relationship, and it was proved that if the total carbon, losson- ignition, and volatiles were kept in the desired ranges, the casting performance was acceptable. If any of the results fell into the grey area (the area of concern) then positive action had to be taken. Equally, going out of the maximum area also required reduction of additives. Sulfur levels normally mirrored the total carbon and ran in the range 0.06 to 0.09.
A testing unit used at Lincoln Castings Ltd. was supplied by Leco UK — a SC144 carbon/sulfur machine. It was selected because it determines carbon/sulfur in various organic matrices from low concentrations (e.g., bentonite) to high levels, such as coal and coke, as well as higher carbon products, such as recarburizers and graphites. This instrument offers a simple solution that determines carbon and sulfur simultaneously using direct combustion and infrared detection.
No hazardous chemicals are used and accurate results are provided in 3 minutes from a small sample weight of 0.3 grams, using reusable refractory sample boats, without accelerating elements. Coupled to the unit is an automatic weighing unit with a dedicated software package, allowing for data storage, statistical analysis and customized operating parameters.
The samples (such as bentonite, coal, silica sand, as well as system and return sand and extraction fines) are tested on a daily basis to ensure proactive steps can be taken, coupled with the routine sand testing. Analysis begins with a nominal weight of 0.3 grams being weighed in a combustion boat. The sample is placed in a pure-oxygen environment regulated at 1,350ºC. The combination of high temperature and oxygen flow cause the sample to combust, and to go through an oxidation and reduction process that causes carbon and sulfur bearing compounds to break down and free the carbon and sulfur. The carbon then oxidizes to form CO2 and the sulfur forms SO2.
The combustion system allows the sample gases to remain in the high-temperature zone, and this permits efficient oxidation. Gases flow through two anhydrone moisture traps and into the infrared detection zone. Then, the carbon and sulfur dioxide values are converted, taking into account the sample weight, calibration, and known moisture value.
The SC 144 can be classed as high-end, multi-element detection instrument, and various laboratories worldwide were checked to determine the effects of alternative units. Many labs have older carbon and carbon/sulfur units, doing mostly metal analysis. It was decided to check samples on both an SC 144 and a CS 244 to determine the difference in results based on a wide range of samples.
The CS 244 and other units in the CS series are ideal for smaller, lower-volume laboratories looking for a cost-effective solution, without sacrificing precision, reliability, or accuracy. These CS units use induction heating elements and a larger sample weight of 1 gram. Also used are accelerating elements, such as high-purity iron, tungsten, or copper with a temperature in excess of 2,000ºC for 15 seconds, to ensure complete combustion. Any CO or SO3 is converted to ensure that only CO2 and SO2 are measured. Both of these gases absorb infrared energy at precise wavelengths within the spectrum. Energy is absorbed as the gases pass through the infrared absorption cells.
The CS unit uses a refractory crucible, and a lid is recommended to prevent any of the sample being removed by the pressure blast of oxygen used in this unit. Research that compared their performances showed the CS cannot measure the initial fast release of volatile gas and therefore the total carbon readings are around 1% lower than when using the SC unit.
Both units are calibrated using a calcium carbonate at 12% carbon and 0.0106% for sulfur. Calibration graphs and running widely varying samples show the units are very accurate, with excellent reproducibility.
As a control measure, many foundries have relied on loss-on-ignition and volatiles tests to measure carbonaceous addition. Coupled with AFS clay grade and active clay measurement, this was considered adequate testing, along with the usual series of permeability, strength, and moisture tests. Experience has shown that active clay levels in most systems range from 1% to 2% higher than required, and these levels are regarded as a safety feature. Whereas we know foundry workers make excellent moldable sand, this does not always translate into top quality castings.
Volatiles testing can be difficult and total carbon determination removes the need for total reliance on this method. A common problem with volatiles testing is the use of a wideneck crucible with a poor fitting lid. These units always give a larger false reading and the use of the standard parallel-sided crucible is urged, with a tight fitting lid. Loss-on-ignition is of course an easier test and this can be backed up with further testing on washed and unwashed sand samples. Conduct both volatiles and loss-on-ignition tests before and after washing out the clay grade, to determine the contribution from the carbonaceous additive. Approximately 1% of volatiles in a sand system actually comes from the bentonite addition.
One foundry does not have all the answers, so it was important to collect sand samples from as wide a group as possible, By selecting samples from foundries Malaysia, Thailand, Czech Republic, Denmark, South Africa, India, as well as the U.K., it was possible to include all available core binders and also differing silica sands, bentonites, and carbonaceous additives, including those containing lustrous carbon.
All participating foundries were assessed for optimum use of raw materials, control of sand testing, and molding properties, to ensure a level playing field. If the locations met the criteria, sand samples were collected over a period of time (all through 2005) to look at trends and importantly at casting performance.
The key measurables were casting scrap related to sand condition, surface-finish quality, sand carryover at knockout, and shotblast times. Of course, not all of this can be attributed to carbonaceous additive but in the end I settled on a good, average, and poor ratings. This may at first seem too simple to be useful, but I believe the results and conclusions justify the selection. Now, further work needs to be done to put numbers against the various selected categories.
Having defined the foundry, samples were collected over the year, and washed and unwashed samples were tested for volatiles and loss-on-ignition, as well as using the SC 144 unit to determine the total carbon and sulfur levels. I personally visited each foundry, in order to monitor casting performance and to crosscheck our own laboratory results. It soon became apparent that the categories matched the level of total carbon within bandwidths.
As no database of information was available, over 20 foundries were tested with total carbon levels ranging from a low 1.5% to a high of 6.68%. Out of interest, and because the SC 144 also measures sulphur at the same time, we recorded the numbers and found a range from 0.045% to 0.13%. Against each sample we conducted volatiles and loss-on-ignition tests to see if there was a relationship, and to check the foundries’ own results (which in itself was a useful exercise): quite differing results were obtained. Three independent laboratories were used, including foundries and suppliers, to verify results over the year. The table shows the results from a single session using various samples on the SC 144 unit, from low- to high-carbon content. The CS results show the difference in the detection systems, but the accuracy shows the possibilities and the versatility of the units for rapid testing and control.
The main list shows the results from the average of samples taken from each foundry’s system sand. Both volatiles and loss-on-ignition tests were carried out, too, to supplement the total carbon results. All these results come from the SC 144 unit, so we may be confident we have measured the total carbon content.
Evaluating volatiles, loss-on-ignition properties The whole matter of the use of carbonaceous additives and their purpose has been well documented. This work simply adds to the conclusion that there is an interaction between the total carbon in a sand system, allied to a level of active effective volatiles, and a loss-on-ignition maximum for successful casting production. Often overlooked in problem solving is the amount of volatiles, which must come from the bentonite, and typically this is classified as 10% of the AFS clay grade. To investigate this further, washed and unwashed tests reveal the true active volatiles content from the carbonaceous additive. Many foundries are simply unwilling to move their loss-on-ignition numbers above 5%, which limits the active volatiles, which in turn limits casting performance.
The graph shows the approximate levels of volatiles for good casting performance from results obtained in one foundry. Each foundry could construct its own graph based on the parameters selected. Obviously other factors come into play but it is clear that the quality active volatiles release is one of the major factors for casting excellence. The “classic” loss-on-ignition (LoI) to volatiles ratio of 3:1 still holds true in most cases, and when further linked to total carbon levels the picture becomes clearer.
Clay-grade LoI is the difference between the washed and unwashed LoI tests. Adding this difference to total carbon will produce the total LoI. Importantly, it shows the correlation between the tests. The example listed shows a clean sand system lacking active volatiles and low in total LoI.
Other control methods The Mass Balance predictive method is used to ensure heat losses after casting (burn out) are reacted to before pouring, and to ensure that additional additives are added or decreased ahead of heavier or lighter castings. This method finds favor in well organized foundries that control additive levels effectively. It allows for lower levels to be added as normal and larger additions ahead of the heavier casting load. This depends on the sand-to-metal ratio, but with all foundries aiming for maximum efficiency it makes good sense to understand how the system will react to pattern changes and varying heat loads.
This system can be incorporated into the sand plant computer controls, and can equally be used for manual adjustments, using what some foundries call the “traffic light” system. They use red to indicate the heaviest castings on the program, amber for mid-range products, and green for lighter heat loads on the system. Irrespective of the system, the important aspect is to be aware of the system, the heat loads proposed, and to be able to take action before corrective measures are required. Rushed changes based on “feel” may work, but informed change is the safer, more positive approach.
Return sand control is one of the most neglected operating practices in green sand foundries. For most foundry workers, what happens to sand after casting is a mystery. What manages to stay in the system simply re-appears for mulling, and so it goes on. Here is a massive area of potential savings, coupled with the knowledge that a well-prepared return sand will be lower in temperature, higher in moisture (ideally at 2%+), and better developed in terms of bentonite and general condition.
A well-developed return sand is much easier to re-mull. In fact, lower addition levels will be required, which in turn leads to lower moisture in the system sand. With lower moisture levels the conditions for a reducing atmosphere are better and, coupled with adequate carbon levels and active volatile content, this will ensure better performance and result in improved casting quality
Casting knockout is another standard operating practice that often is overlooked as an area of potential quality control. The results of all calculations and control are available simply by looking critically at the resultant castings. The sand peel from the castings, the amount of sand carried over, and the color of the castings is often a good early indicator of casting quality.
Simple, complex, effective Green sand systems can be simple and complex — simple in terms of additives and control; complex in the amount of variables possible and the need for constant vigilance — but they are simply the most cost-effective method for highvolume production of iron castings.The research shows that total carbon can be a very useful measuring and control tool for green sand systems, coupled with accurate volatile and loss-on-ignition tests. Without exception the foundries categorized as acceptable in this research had total carbon figures in excess of 3%. In foundries categorized as good, it was no coincident that the volatile and losson- ignition figures were also in what the author describes as “a good area.”Foundries would do better to arm themselves with meaningful data regarding carbonaceous additives. Total carbon determination could include all aspects of a sand system, including raw materials, and this would simplify decision-making. Areas such as return sand and fines extraction could be monitored easily with better control.All sand systems have differing moisture sensitivity. Using bentonite from known, approved sources coupled with correctly graded coal, with low ash content (less than 3%), ensures the coal does not compete with the bentonite for available water. The key for foundries is to understand the properties of raw materials they use, and to control the consistency. Once this has been established, it will be possible to rely on a minimum amount of additives to achieve the goal of consistent quality castings. Ensuring adequate carbonaceous levels is not difficult but for many foundries it remains only a secondary focus. World-class performance is being demanded by casting buyers, and successful suppliers will need to ensure they operate “in the correct zone,” with effective knownsource additives.

aluminium in cast iron

2009-08-06T00:33:00.000-07:00

Aluminium in Cast Iron
Aluminium is normally found in cast irons as a mainly harmless residual element. The
major sources for aluminium are steel scrap, contaminated cast scrap (engine blocks etc
with pistons included), the ferroalloys consumed and inclusions of non-ferrous metals in
the charge materials.
A common occurrence in foundries is the pinhole problem from hydrogen gas evolution,
which often can be attributed to excessive aluminium contents. It is accepted that alumi-
nium has an influence on the surface tension of the liquid iron, a consequence of which
could be susceptibility to pinholing defects. The figure below shows the relationship bet-
ween aluminium in the iron and the tendency to pinholing. It is shown that grey iron is
more sensitive to pinholing than ductile iron due to the overall lower surface tension for
grey iron. Above a certain level (approximately 0.2% Al), the susceptibility for pinholing is
reduced as the surface tension again increases. The most critical range is 0.05 – 0.2% for
ductile iron and 0.008 – 0.2% for grey iron. Consequently, the contents of aluminium
should always be kept low, preferably below this range where the risk will be highest.
Influence of Aluminium Content on Surface Tension and Pinhole
Susceptibility of Grey and Ductile Irons
It should also be noted that iron temperature will influence the surface tension and thus
well insulated ladles are of importance (refer to Elkem Technical Information Sheet No. 21
for more data). Aluminium will also add to the slag formation, resulting in poor furnace per-
formance, more ladle and holder maintenance, and increased risk for slag inclusions in
castings.
Aluminium has virtually no inoculating effect as such, but it may add to the hardness of the
iron, and it can also be harmful to the nodularity of ductile iron. It is also important to note
that titanium will play the same role as aluminium to a certain extent, although normally
present in smaller amounts than aluminium. Furthermore, the two elements will have an
aggregated effect, and both elements should be monitored and controlled at all times.
Elkem ASA, Foundry Products
Postal address
P.O.Box 5211 Majorstuen
NO-0302 Oslo
Norway
Office address
Hoffsveien 65B
Oslo
Norway
Telephone
+47 22 45 01 00
Telefax
+47 22 45 01 52
© Copyright Elkem ASA
Web
www.foundry.elkem.com
Org. no.
NO 911 382 008 MVA
Revision
No. 1.1
20.03.2004
Technical Information 19
2
It is also well known that many elements can interact with aluminium to affect the iron pro-
perties, either by enhanced inoculation potency or by detrimental effects such as the com-
bination of aluminium and titanium. The presence of even minor traces of titanium means
that the tolerable aluminium levels will be dramatically reduced. The figure below shows
the combined effects of aluminium and titanium on the hydrogen pinholing tendency in
ductile iron. Above the curve there will be a significant risk for such defects to occur.
Combined effects of aluminium and titanium on
hydrogen pinholing tendency in ductile iron.
Example of pinhole defect in grey iron.
For ductile iron the permissible aluminium is roughly 5 – 10 times that of grey iron. No data
is available concerning the combined effects of aluminium and titanium in grey iron, but
there are reasons to believe the interaction is about as for ductile iron and both elements
should therefore be closely watched.
Since both nodularizers (magnesium-ferrosilicon) and inoculants will contain various
amounts of aluminium and titanium, it is important that choice of alloys is being made with
full awareness of its total chemical composition. At higher aluminium contents, ferrosilicon-
based alloys will tend to improve solubility, but the increased slag formation and pinholing
tendency should call for caution. High aluminium containing alloys should hence only be
used where low addition rates are applicable (i.e. stream inoculation). Special attention
should be paid to large amounts of ferrosilicon used as a furnace charge material.
It is worth noting that hydrogen pinhole defects often will have similar characteristics as
other type of gas defects, such as nitrogen porosity. A characteristic feature for pinholes is
the graphite lining covering the inner pore surfaces. An example of such a hydrogen
pinhole defect is shown in the figure above. This defect characteristic can also occur for
nitrogen defects, and it is therefore often difficult to separate between such gas defects. A
thorough investigation into nitrogen, aluminium and titanium levels will be necessary to
determine the type of gas involved, since a high Aluminium and Titanium level may
promote hydrogen pinholes but at the same time effectively neutralize nitrogen by forming
TiN and AlN inclusions.
Choice of core binder system and green sand humidity level is also vital for the avoidance
of hydrogen and nitrogen pinhole defects.

solidification of cast iron

2009-08-06T00:30:00.000-07:00

در پيام بعدي برايتان ارسال خواهم كرد

ايميل و تلفن هاي من

2009-08-05T23:28:00.000-07:00

لطفا مسائل ريخته گري خودتان را با ما در جريان قرار دهيد

09144077029

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مقاله جديد

2009-08-05T22:57:00.000-07:00

در پيام بعدي مقالات زير را براي استفاده در اختيارتان قرار خواهم داد
- اصلاح الگوی مصرف در ریخته‌گری
- Mold inoculation of cast iron using pressed blocks
- core coating

in mold inoculation

2009-08-05T22:55:00.000-07:00

بررسي تاثير عوامل مهم در فرآيند توليد چدن داكتيل به روش افزودن مواد در داخل قالب[1]
نگارش:واحد نجاتي مازگر
شركت ريخته گري تراكتورسازي ايران
خلاصه:
در اين تحقيق اثر عواملي چون درصد وزني و اندازه ذرات مواد كروي كننده فرو سيليكو منيزيم ،طرح سيستم راهگاهي و ابعاد محفظه واكنش در فرآيند توليد چدن داكتيل به روش افزودن مواد كروي كننده در داخل قالب(اينمولد) خواهيم پرداخت.همانطور كه مي دانيم با بكارگيري فرآيند اينمولد از ميرائي منيزيم جلوگيري به عمل مي آيد چرا كه مذاب تلقيح شده بلافاصله وارد محفظه قالب مي گردد.نيز از آنجا كه واكنش منيزيم با مذاب در محيط بسته عاري از اكسيژن رخ مي دهد ميزان سوختن منيزيم نيز به شكل قابل ملاحظه اي كاهش يافته و بازيابي منيزيم تا 80 درصد افزايش مي يابد.در كنار اين مزايا و با در نظر گرفتن عدم مسائل زيست محيطي و بالا بودن ايمني كار همراه با صرفه هاي اقتصادي و سادگي روش تمايل به استفاده از اين روش نزد ريخته گران به شدت افزايش يافته است .لذا با توجه به استفاده گسترده ريخته گران از اين روش در توليد قطعات داكتيل و نيز با عنايت به كمبود اطلاعات مدون در رابطه با پارامترهاي مهم و موثر اشاره شده ، در تحقيق حاضر بر آن شديم تاعوامل موثر مذكور را مورد تجزيه و تحليل قرار دهيم و انشااله كه نتايج حاصله مورد استفاده ريخته گران محترم قرار گيرد.
مقدمه:
استفاده از روش اينمولد براي تلقيح مواد جوانه زا و نيز كروي كننده در توليد چدن داكتيل امروزه جايگاه ويژه اي را نزد ريخته گران دارد.اولين بار در اواسط دهه 60 از اين روش جهت جوانه زني چدن داكتيل (با افزودن فروسيليسيم 85 درصد در مسيرسيستم راهگاهي قطعه) صورت گرفت.پس از آن در ابتداي ده 70 استفاده از اين روش در اروپا و آمريكا نيز رونق يافت و آنها از مواد جوانه زا به شكل بلوكه هاي ريخته يا پرسي سينتر شده ،در داخل قالب استفاده به عمل آوردند.در نهايت در اواسط دهه 70 اين روش براي تلقيح مواد كروي كننده در توليد چدن داكتيل مطرح گرديد.از آن زمان تا بحال اين روش دائما دستخوش تغييرات و بهينه سازي بوده و امروزه به عنوان فراگيرترين روش در ريخته گري چدن داكتيل استفاده مي گردد. در كنار مزاياي اشاره شده براي اينمولد، اين روش نيز همانند ديگر روش ها داراي معايبي مي باشد و عدم كنترل برخي پارامترها باعث ضايعات جبران ناپذير قطعات مي گردد كه در ساير روشها اين حساسيت ها كمتر مي باشد.از جمله مهم ترين معايب اين روش مي توان به كاهش بهره ريختگي (به علت مذاب اضافي در محفظه واكنش)، بالاتر بودن احتمال ورود آخال به داخل قطعه و حساسيت بالاي اين روش به مشخصه هاي ماده كروي كننده مورد استفاده اشاره نمود.حال با طراحي صحيح سيستم راهگاهي و محفظه واكنش در كنار درك صحيح از اثرات مربوط به ماده كروي كننده ،براحتي مي توان محدوديتهاي اشاره شده را به حداقل رساند.در تحقيق حاضر آزمايشات متعددي در اين زمينه ها انجام و به شرح زير مورد بحث و بررسي قرار گرفت.
روش انجام تحقيق:
مراحل انجام اين تحقيق را مي توان به دو بخش به شرح زير تقسيم بندي نمود:
مرحله اول شامل تعيين مقدار و اندازه بهينه آلياژ كروي كننده و
مرحله دوم مربوط به طراحي سيستم راهگاهي و محفظه واكنش در روش اينمولد.
طراحي آزمايشات مربوط به مرحله اول:
براي انجام آزمايشات اين مرحله مذاب گوگرد زدائي شده با كاربيد كلسيم در كوره القائي 14 تني با فركانس بالا تهيه گرديد.جهت كنترل آناليز ابتدا نمونه پولكي ازجنس چدن سفيد با ريختن مذاب كم گوگرد در داخل قالب مسي و سرد كردن سريع آن با آب تهيه شد و پس از سنگزني با استفاده از دستگاه كوانتومتري با مشخصات .......... آناليز آن تعيين گرديد .در جدول 1 آناليز نمونه هاي ريخته شده در اين تحقيق قبل ازتلقيح منيزيم آورده شده است.آلياژ FeSiMg به ميزان 1.25,1,0.75,0.5 و1.5 درصد وزني مذاب با اندازه ذرات به ترتيب 1mm,4mm,6mm و زير 1mm در هر كدام از اوزان مذكور در داخل محفظه واكنش در دو حالت به شرح زير ريخته شدند.
الف) تنگه و راهبار با محفظه واكنش در امتداد يك خط راست باشند(شكل1b ).
ب) تنگه و راهبار با محفظه واكنش در امتداد يك خط راست نباشند(شكل1a ).
سيستم قالبگيري مورد استفاده در اين تحقيق به روش ماسه اي تر مي بود كه خواص مخلوط ماسه نيز در جدول 2 آورده شده است.

جدول1)خواص مخلوط ماسه قالبگيري مورد استفاده در اين تحقيق
تراكم(%)
رطوبت(%)
استحكام كششي(MPa )
پودر زغال(%)
بنتونيت(%)
عدد ريزي(AFS)
43-40
3.7-3.4
18-16
3.6-3.4
11-9
52

جدول2) تركيب شيميائي نمونه هاي ريخته شده در اين تحقيق
كربن معادل
كربن
سيليسيم
گوگرد
منگنز
منيزيم
فسفر
4.3
3.7-3.6
2.2-2
0.013≥
0.35-0.25
0.05-0.045
0.03

روش انجام آزمايشات مرحله اول:
با توجه به آنچه در شكل (a,b) 1 نشان داده شده است ،سيستم راهگاهي در اين تحقيق شامل راهبار،راهباره ها ،حوضچه بارريزي،حوضچه پاي راهگاه،راهگاه بارريز،تنگه و محفظه واكنش مي باشد كه مذاب را به محفظه Y– بلوك با ابعاد استاندارد هدايت مي كند. تنگه در اين سيستم در ورودي و خروجي محفظه واكنش قرار داده شد و نيز براي جلوگيري از ورود آخال به محفظه قالب راهبار در قالب بالائي تعبيه گرديد.بررسي هاي انجام پذيرفته نشان مي دهد زماني كه تنگه و راهبار با محفظه واكنش در امتداد يك خط راست قرار گرفته باشند(شكل1b ) در تمامي شرايط آزمايش واكنش بين مواد كروي كننده با مذاب كامل نمي باشد و پس از برش محفظه واكنش مقداري مواد واكنش نيافته در انتهاي محفظه باقي مي ماند كه آنهم باعث افت ندولاريته در تمامي نمونه ها مي گردد و لذا تمامي آزمايشات با سيستم راهگاهي حالت دوم كه تنگه و راهبار با محفظه واكنش در امتداد يك خط راست نمي باشند)شكل(1a، انجام پذيرفت.
در مرحله بعد مواد كروي كننده به ميزان 0.5%,0.75%,1.0%, 1.25%,1.5% با دانه بندي 1,4,6mm و زير 1mm براي هر كدام از مقادير وزني ذكر شده در محفظه واكنش ريخته شد(در واقع 20 عدد Y– بلوك با شرايط فوق توليد گرديد).دماي ذوبريزي در تمامي موارد 1420±5ºС در نظر گرفته شد.از Y– بلوكهاي ريخته شده نمونه متالوگرافي تهيه و بررسي هاي ريزساختاري روي آنها به عمل آمد.

راهباره
راهبار

(a)
Y - بلوك
محفظه واكنش
حوضچه
(b)

شكل1) سيستم هاي راهگاهي مورد استفاده در دو حالت (a) :راهبار و تنگه با محفظه واكنش
در امتداد يك خط نيستند و (b) : راهبار و تنگه با محفظه واكنش در امتداد يك خط هستند

بحث و نتيجه گيري:
نتايج حاصل از با در نظر گرفتن اندازه ذرات 6 ميليمتر در مقادير مختلف وزني در جدول (3) نشان داده شده است.همانگونه كه ديده مي شود در 0.5 و0.75 درصد وزني آلياژ كروي كننده ساختار گرافيت لايه اي ظاهر مي گردد.با افزايش مقدار مواد تا 1.25 درصد ،تعداد و درصد گرافيتهاي كروي نيز افزايش مي يابد .در نهايت با افزايش مقدار مواد كروي كننده به 1.5 درصد ندول كانت مجددا كاهش مي يابد.تصاوير ريزساختاري مربوط به اين نمونه ها در شكل 2 نشان داده شده است.
جدول 3 ) نتايج حاصل از بكارگيري ذرات FeSiMg با اندازه 6 ميلي متر
1.5%
1.25%
1%
0.75%
0.5%
درصد آلياژ FeSiMg
90
110
48
-
-
ندول كانت
95
95
60
لايه اي
لايه اي
ندولاريته (%)
e
d
c
b
a
عنوان شكل

c b a

e d
شكل2) ريزساختار نمونه هاي ريخته شده با شرايط مندرج در جدول 3

در مرحله بعدي اندازه ذرات مواد كروي كننده به 4 ميليمتر كاهش داده شد.همانگونه كه در ريزساختارهاي حاصله نشان داده شده است در اين شرايط حتي با 0.5 درصد مواد كروي كنند ه نيز 35 درصد گرافيتها به شكل كروي در مي آيد و ريزساختار حاوي گرافيتهاي فشرده و كرمي شكل ظاهر مي گردد.به تدريج با افزايش مقدار مواد ندول كانت و درصد گرافيتهاي كروي نيز افزايش مي يابد و برخلاف حالت قبل هيچ افتي در ميزان ندول كانت با بكارگيري 1.5 درصد مواد كروي كننده بوجود نمي آيد.
جدول 4 ) نتايج حاصل از بكارگيري ذرات FeSiMg با اندازه 4 ميلي متر
1.5%
1.25%
1%
0.75%
0.5%
درصد آلياژ FeSiMg
109
175
90
50
-
ندول كانت
95
95
85
65
لايه اي
ندولاريته
e
d
c
b
a
عنوان شكل

c b a

e
d
شكل3) ريزساختار نمونه هاي ريخته شده با شرايط مندرج در جدول 4

در مرحله بعدي اندازه ذرات مواد كروي كننده به 1 ميليمتر كاهش داده شد.همانگونه كه در جدول 5 و ريزساختارهاي مربوطه در شكل 4 نشان داده شده است با كاهش اندازه ذرات فقط در مقادير 1.25 و 1.5 درصد وزني مواد كروي كنند ه ريزساختار حاوي مقداري گرافيت كروي بدست مي آيد.در ادامه آزمايشات ملاحظه گرديد كه با كاهش اندازه ذرات به 1 ميليمتر و كمتر،در هيچ كدام از درصد هاي وزني ساختار حاوي گرافيتهاي كروي بدست نمي آيد و ريزساختار كلا خاكستري حاصل مي گردد.نتايج حاصل از اين آزمايشات نيز در جداول و اشكال زير نشان داده شده است.

جدول 5 ) نتايج حاصل از بكارگيري ذرات FeSiMg با اندازه 1 ميلي متر
1.5%
1.25%
1%
0.75%
0.5%
درصد آلياژ FeSiMg
87
64
5
-
-
ندول كانت
90
75
10
لايه اي
لايه اي
ندولاريته
e
d
c
b
a
عنوان شكل

c b a

e d
شكل4) ريزساختار نمونه هاي ريخته شده با شرايط مندرج در جدول 5

در مرحله بعدي اندازه ذرات مواد كروي كننده به زير 1 ميليمتر كاهش داده شد.همانگونه كه در جدول 5 و ريزساختارهاي مربوطه در شكل 4 نشان داده شده است با كاهش اندازه ذرات فقط در مقادير 1.25 و 1.5 درصد وزني مواد كروي كنند ه ريزساختار حاوي مقداري گرافيت كروي بدست مي آيد.در ادامه آزمايشات ملاحظه گرديد كه با كاهش اندازه ذرات به 1 ميليمتر و كمتر،در هيچ كدام از درصد هاي وزني ساختار حاوي گرافيتهاي كروي بدست نمي آيد و ريزساختار كلا خاكستري حاصل مي گردد.نتايج حاصل از اين آزمايشات نيز در جداول و اشكال زير نشان داده شده است.

جدول 6 ) نتايج حاصل از بكارگيري ذرات FeSiMg با اندازه 1 ميلي متر و كمتر
1.5%
1.25%
1%
0.75%
0.5%
درصد آلياژ FeSiMg
5
-
-
-
-
ندول كانت
15
-
-
لايه اي
لايه اي
ندولاريته
e
d
c
b
a
عنوان شكل

c b a
e d
شكل5) ريزساختار نمونه هاي ريخته شده با شرايط مندرج در جدول 6

بحث و نتيجه گيري از آزمايشات مرحله اول:
نتايج بدست آمده به وضوح نشان مي دهد بهترين اندازه براي ذرات كروي كننده FeSiMg و ميزان افزودن آن به محفظه واكنش4 -2 ميليمتر به مقدار 1.25-0.9 درصد بسته به درصد منيزيم محتوي آلياژ مي باشد .حال با ثابت در نظر گرفتن اين پارامترها نسبت به تبيين پارامترهاي طراحي اقدام گرديد.
طراحي آزمايشات مربوط به مرحله دوم:
در كنار جنس ،اندازه و مقدار مواد كروي كننده مسائل مربوط به طراحي سيستم راهگاهي و محفظه واكنش نيز از اهميت بالائي برخوردار مي باشد.حال با توجه به اينكه در آزمايشات مرحله اول از سيستم هاي ساده و متداول در طراحي روش اينمولد استفاده بعمل آمد در اين مرحله خواهيم ديد كه طراحي صحيح آنها نقش به سزائي در حصول ريزساختار و خواص مورد نظر از چدن داكتيل را بر عهده دارند.
همانطور كه مي دانيم مهم ترين پارامتر در طراحي محفظه واكنش سطح مقطع افقي ،عمق و شيب ديواره آن مي باشد.براي تعيين اين پارامترها نتايج تحقيقات ساير محققان نيزمد نظر قرار گرفت.ولي با توجه به حساسيت بيش از حد روش اينمولد به متغير هاي متعدد به هيچ عنوان نمي توان عين نتايج حاصل از يك تحقيق را در توليد با شرايط متفاوت به كار گرفت و فقط در كم كردن تعداد آزمايشات مي تواند موثر واقع گردد.لذا براي جلوگيري از تكرر آزمايشات و نيزبا توجه به اينكه شركت ELKEM از جمله پيشتازان توليد كننده مواد براي روش اينمولد مي باشد ،سيستم راهگاهي پيشنهادي از سوي محققان اين شركت را در كليه آزمايشات لحاظ كرديم و در ادامه آزمايشات براي تعيين ابعاد محفظه واكنش صورت پذيرفت.در شكل 6 نحوه محاسبه مقاطع سيستم راهگاهي نشان داده شده است.
شكل6) طرح شماتيك پيشنهادي از طرف شركت ELKEM براي روش Inmold و محاسبات مربوطه

براي محاسبه سطح مقطع محفظه واكنش از رابطه زير استفاده مي گردد:
Chamber Area = Pouring Rate/ Alloy Solution Factor
حال با توجه به اينكه نرخ ذوبريزي در حين توليد مقدار ثابتي مي باشد لذا مهم ترين پارامتر تاثيرگذار در ابعاد محفظه واكنش ضريب حلاليت آلياژ كروي كننده مي باشد.براي تاييد اهميت موضوع در تحقيق حاضر ابتدا براي يك آلياژ معين و در ضريب حلاليت ثابت ابعاد محفظه واكنش را محاسبه كرده و سپس آلياژهاي با ضرايب مختلف را در همان محفظه طراحي شده آزمايش كرديم كه نتايج حاصله گواه اين مدعي مي بود كه ابعاد محفظه واكنش كاملا وابسته به ضريب حلاليت آلياژ بوده و لحاظ نكردن دقيق آن در محاسبات باعث بروز مشكلات ريزساختاري و در نهايت خواص نهائي قطعه مي گردد.
روش انجام آزمايشات مرحله دوم:
با توجه به اطلاعات ارسالي از طرف فروشنده ضريب حلاليت مواد كروي كننده FeSiMg مورد استفاده در اين تحقيق 0.07[kg/cm2sec] مي بود.زمان ذوبريزي در هر قالب 8 ثانيه و وزن كل مذاب ريخته شده به قالب تقريبا 32 كيلوگرم مي بود .لذا نرخ ذوبريزي و سطح مقطع محفظه واكنش در اين آزمايشات به شرح زير محاسبه گرديد:
Pouring rate= W / t =32 / 8 = 4 kg/sec
Chamber Area= 4 / 0.07 =57.14 cm^2 :و از روي معادله فوق
اين عدد در واقع متوسط سطع مقطع افقي محفظه واكنش مي باشد.حال با توجه به دانسيته مواد (2.16 gr/cm^3) و وزن مورد استفاده (1.1 درصد وزني مذاب ريخته) مواد، حجم آن به شرح زير خواهد بود:
=دانسيته / وزن=حجم مواد(محفظه) 32×1000×1.1 / 100×2.16 =163 cm^3
لذا ارتفاع محفظه برابر خواهد بود با: cm 163 / 57.14 = 2.85 = ارتفاع محفظه
در طراحي عموما يك اينچ فضاي خالي روي ارتفاع محاسبه شده لحاظ مي گردد و لذا ارتفاع محفظه واكنش تقريبا 5.4 سانتي متر و با در نظر گرفتن 6 ميلي متر ارتفاع كانال ورودي ذوب به محفظه عمق نهائي محفظه واكنش 6 سانتي متر لحاظ گرديد و مقطع متوسط آن 57 سانتي متر مربع در نظر گرفته شد. همچنين در جدول 7 و شكل 7 خواص و ريز ساختار قطعات ريخته شده با اين محاسبات نشان داده شده است.

جدول 7 ) نتايج حاصل از بكارگيري ذرات FeSiMg با اندازه 4 ميلي متر با سيستم راهگاهي
حاصل از محاسبات فوق و نشان داده شده در شكل 7
زمينه ريزساختار
ندولاريته
ندول كانت
درصد آلياژ FeSiMg
Fe+15%Pe
95
175
1.1%

شكل7) ريزساختار نمونه هاي ريخته شده با شرايط مندرج در جدول 7

در مرحله بعدي آزمايشات نرخ ذوبريزي با تغيير دادن سطح مقطع راهگاه بارريز و ورودي ذوب به Y – بلوك تغيير داده شد.سپس با استفاده از محفظه واكنش طراحي شده در مرحله قبل با ضرايب حلاليت مختلف بدست آمده آزمايشات به شرح زيرادامه يافت.هدف از آزمايشات اين مرحله نشان دادن اهميت بالاي ضريب حلاليت در محاسبات مربوط به محفظه واكنش مي باشد.
در اولين آزمايش از اين مرحله نرخ ذوبريزي به kg/sec6 افزايش داده شد.همانطور كه بيان شد علي رغم تغيير در نرخ ذوبريزي و به تبع آن ضريب حلاليت،محفظه واكنش ثابت در نظر گرفته شده است لذا با لحاظ نمودن 57 سانتي متر مربع براي سطح مقطع محفظه ،ضريب حلاليت آلياژ برابر با [kg/cm2sec] 0.105 خواهد بود.ريزساختار مربوطه به نمونه ريخته شده با اين شرايط در شكل8 نشان داده شده است.همانطور كه ملاحظه مي گردد علي رغم بالا بودن ندولاريته برخي قسمتهاي قطعه عاري از گرافيت كروي مي باشد وساختار گرافيت لايه اي ظاهر شده است.علت اين موضوع تمام شدن مواد كروي كننده قبل از اتمام ذوبريزي مي باشد كه باعث شده است از داكتيل شدن قسمت انتهائي قطعه جلوگيري نمايد.

a b
شكل8) ريزساختار نمونه هاي ريخته شده با افزايش نرخ ذوبريزي و ثابت ماندن ابعاد
محفظه در(a):مناطق نزديك به راهباره و (b) :مناطق حاوي آخرين مذاب در Y - بلوك

در آزمايش بعدي نرخ ذوبريزي به kg/sec2 كاهش داده شد و مشابه حالت قبل از همان محفظه واكنش براي توليد نمونه استفاده به عمل آمد.در اين شرايط با كاهش نرخ ذوبريزي ضريب حلاليت آلياژ به شرح زير خواهد بود:
[kg/cm2sec] 0.035 = 57 / 2 = ضريب حلاليت
ريزساختار مربوط به نمونه ريخته شده با اين شرايط در شكل 9 نشان داده شده است.همانطور كه ملاحظه مي گردد درصد ندولاريته به شدت كاهش يافته است و علت آن نرخ پائين حلاليت آلياژ نسبت به سطح مقطع طراحي شده مي باشد.لذا در چنين شرايطي لازم بود سطح مقطع محفظه افزايش داده شود.
شكل9) ريزساختار نمونه هاي ريخته شده با كاهش نرخ ذوبريزي و ثابت ماندن ابعاد محفظه

بحث و نتيجه گيري از آزمايشات مرحله دوم:
همانطور كه ملاحظه گرديد عدم طراحي صحيح محفظه واكنش متناسب با ضريب حلاليت آلياژ ريزساختار حاوي گرافيتهاي غير كروي در مقادير متنابهي در ريزساختار ظاهر شده و بالطبع باعث افت استحكام قطعه مي گردد.در اين رابطه ملاحظه شد كه تغيير در نرخ ذوبريزي اثر مستقيمي روي ضريب حلاليت آلياژ مي گذارد و لذا در چنين شرايطي لازم است ابعاد محفظه واكنش به نسبت آن تغيير يابد.شايان ذكر است تغيير در ضريب حلاليت آلياژ علاوه بر نرخ ذوبريزي مي تواند ناشي از تغييرات دماي ذوبريزي،درصد Mg آلياژ كروي كننده و درصد گوگرد مذاب پايه باشد.همين قضيه حساسيت بالاي روش اينمولد را كه يكي از معايب اين روش مي باشد و در ابتداي مقاله توضيح داده شد ،نشان مي دهد.
نتايج:
در اين قسمت لازم به ذكر است كه در توليد انبوه به روش اينمولد و حتي در طول انجام آزمايشات اين تحقيق يكسري نتايج مغاير با آنچه پيش بيني كرده بوديم بدست آمد كه ذكر علل تمامي آنها از حوصله اين مقاله خارج مي باشد ولي سعي كرديم در نتايج آورده شده به شرح زير تمامي اين موارد را تا حد امكان توضيح دهيم. در مجموع نتايج كلي حاصل از تحقيق انجام پذيرفته را مي توان به شرح زير بيان نمود:
ü بهترين دامنه ابعادي براي ذرات فرو سيليكو منيزيم در روش اينمولد 4-1 ميلي متر مي باشد و انحراف از اين مقدار باعث بروز مشكلات ريزساختاري مي گردد.اگر سايز ذرات بيش از حد بزرگتر باشد باعث افت ندولاريته مي گردد و لازم است در چنين شرايطي مقدار مواد كروي كننده افزايش داده شود.در اين حالت نيز ندول كنت به تدريج كاهش يافته و امكان بروز آخال در قطعه و نيز ريز مك هاي انقباضي در قطعات افزايش مي يابد.اگر اندازه تمامي ذرات از يك ميلي متر كمتر باشد امكان توليد چدن داكتيل وجود نخواهد داشت.
ü بهترين دامنه افزودن فرو سيليكو منيزيم 1.25-1 درصد وزني مذاب ريخته در قالب مي باشد.بديهي كاهش آن باعث افت ندولاريته و افزايش آن همانند حالت قبل امكان بروز آخال و حفرات انقباضي را افزايش مي دهد و نيز باعث بروز تخلخل[2] در سطح قطعه هم مي گردد.
ü براي انحلال يكنواخت مواد كروي كننده و تغيير نرخ حل شدن مواد متناسب با نرخ ذوبريزي لازم است تنگه در خروجي محفظه واكنش اعمال گردد تا محفظه هم همواره پر از مذاب باشد و نيز ضروري است تنگه و راهبار با محفظه واكنش در امتداد يك خط قرار نگيرد.
ü يكنواختي ريزساختاري با استفاده از محفظه واكنش به شكل چهار وجهي بيشتر از محفظه هاي با شكل استوانه اي مي باشد.
ü فاصله دورترين راهباره از محفظه واكنش بين 10-5 سانتي متر لحاظ گردد تا بهتري نرخ انحلال بدست آيد
ü ضريب حلاليت آلياژ نقشتعيين كننده اي در طراحي محفظه واكنش دارد.بهترين دامنه براي ضريب حلاليت در روش اينمولد 0.08-1[kg/cm2sec] مي باشد و از روي آن و با توجه به نرخ ذوبريزي بايستي ابعاد محفظه محاسبه گردد.اگر ضريب حلاليت به هر دليلي از قبيل تغيير در ميزان منيزيم محتوي آليژ،درصد گوگرد مذاب پايه و يا دماي ذوبريزي تغيير يابد بايستي ابعاد محفظه واكنش نيز اصلاح گردد.اگر ضريب حلاليت افزايش يابد لازم است سطح مقطع محفظه واكنش كاهش يابد تا نرخ انحلال آلياژ متناست با زمان ذوبريزي باشد و در غير اين صورت بخشي از قطعه كه حاوي آخرين مذاب مي باشد عاري از گرافيتهاي كروي خواهد بود و اگر ضريب انحلال كاهش يابد بايستي سطح مقطع محفظه كاهش يابد و در غير اين صورت منيزيم حل نشده در كف محفظه باقي خواهد ماند و باعث افت ندولارينته در قطعه خواهد شد.
ü براي كم كردن احتمال برو‍ز آخال لازم است كانال ورودي مذاب به محفظه در نيمه ژائين قالب و كانال خروجي از آن در نيمه بالائي قالب تعبيه گردد.سطح مقطع اين كانال ها هم از اهميت بالائي برخوردار مي باشد و لذا بايستي ژس از طراحي صحيح از بروز ماسه كني در مقاطع آنها جلوگيري به عمل آيد.
ü بروز ماسه كني باعث افزايش حجم مذاب در تناسب با مقدار منيزيم مصرفي مي گردد و مي تواند باعث افت ندولاريته در قطعات توليدي گردد.لذا ماسه كني در روش اينمولد برخلاف ديگر روش ها از اهميت بالائي برخوردار مي باشد.
ü در پاره اي از موارد بيرون زدن ذوب از سطح جدايش درجه ها به حدي كم مي باشد كه پس از چند ثانيه بيرون زدن قطع مي گردد.در چنين شرايطي برخلاف ريخته گري چدن خاكستري و نيز روش هاي پاتيلي براي توليد چدن داكتيل بايستي از ادامه دادن ذوبريزي اجتناب نمائيم چرا كه با ادامه دادن آن ميزان مذاب مصرفي نسبت به منيزيم محاسبه شده داخل قالب افزايش يافته و باعث افت ندولاريته در قطعات خواهد شد.
ü با توجه به اينكه سطح مقطع افقي محفظه واكنش در نرخ حل شدن آلياژ فرو سيليكو منيزيم نقش مهمي را ايفا مي نمايد لذا لازم است ژس از ريختن مواد در داخل محفظه سطح آن صاف گردد چرا كه در حالت تحدب سطح در تماس با مذاب افزايش در ابتداي ذوبريزي افزايش يافته و با افزايش نرخ حل شدن ممكن است قبل از اتمام ذوبريزي آلياژ تمام شده باشد كه آن هم باعث شكل گيري گرافيت لايه اي در مناطق حاوي آخرين مذاب خواهد شد.
ü پيش گرم كردن آلياژ تا C°160 قبل از ريختن آن در محفظه واكنش لايه هاي اكسيدي و رطوبت احتمالي روي مواد را از بين برده و باعث بهبود كيفي قطعات مي گردد. لازم به ذكر است زمان نگهداري آلياژ،ميزان رطوبت محيط و طراحي هوپر نگهداري مواد روي ميزان اكسيداسيون مواد نقش اساسي را دارند.

منابع و مراجع:
§ Elkem, Technical Information 35
§ Common Metallurgical Defects in Grey Cast Irons,
§ Augus H T. Cast Iron. Butter Worths Co., 1976: 172.
§ Chisamera, M., Riposan, I. – Romanian Patent No.95564, 1986
§ Csonka, J., M, Muratore, E.C. and Woods, J.E. – “Survey on Ductile Iron Practice”, AFS Transactions, Vol. 110, pp.1099-1112 (2002)
§ Naro, R. L. , Wallace, J. F., "Trace Elements in Cast Irons", Trans AFS, Vol. 77., p. 311, 1969
§ Suarez, O.M., Kendrick, R.D. and Loper, C.R.Jr. “Late Sulfur Inoculation of Spheroidal Graphite Cast Irons”. SCI-7, Int. Cast Irons Conference, Barcelona-Spain, Sept.2002.

تشكر و قدرداني:
در نهايت از زحمات تمامي همكاران عزيز و زحمتكش شركت ريخته گري و بخصوص كارگاه ذوب و امور كنترل كيفي و جناب آقاي مهندس صباحي (بخاطر راهنمائي هاي ارزشمندشان) كه در انجام اين تحقيق مرا ياري فرمودند كمال تشكر را داشته و از جناب آقاي مهندس مهاجراني كه بدون حمايت بي دريغ ايشان از فعاليتهاي تحقيقاتي در شركت ريخته گري تراكتورسازي امكان اين تحقيق فراهم نمي شد كمال امتنان را دارم.

[1][1] -In Mold Treatment
[2] -Pitting

Ductile Iron Inoculation

2009-08-01T00:16:00.000-07:00

A New Approach to Ductile Iron Inoculation
T. Skaland
Elkem ASA, Research
Kristiansand, Norway
Copyright©2001 American Foundry Society
ABSTRACT
The objective of the present paper is to describe a new approach to inoculant design that has proven successful in improving
casting performance and properties. The described inoculant represents a unique new generation of products developed for
powerful cast iron inoculation. The ferrosilicon alloy contains levels of Calcium and Cerium that are adjusted to minimize
chill formation and neutralize subversive trace elements in the iron. In addition, the inoculant contains small and controlled
amounts of Sulphur and Oxygen in a form that make them available for reaction with the Calcium and Cerium during
introduction into liquid iron. This special composition is designed to give highly powerful graphite nucleation conditions in
ductile irons along with very effective chill and shrinkage reduction. Examples from foundry testing are reviewed, and the
unique characteristics of this new inoculant concept described. The new inoculant is found to be more powerful than
conventional ferrosilicon based inoculants, and give rise to very effective reduction in the shrinkage tendency of ductile irons.
Special effectiveness has been observed in irons of low sulphur or irons of a “dead” nature from prolonged holding times.
Also, results show improvements in both tensile properties as well as machinability for ferritic ductile irons. The new
inoculant concept represents a patent protected design (Int.Pat.No.WO99/29911), and is available under a special trade name,
see footnote1 .
INTRODUCTION
Based on years of laboratory work with experimental inoculants of various compositions, a new concept for inoculation of
ductile iron has been developed. The concept is novel in the sense that it involves not only an alloyed ferrosilicon-based
material, but also introduction of non-metallic powders with the ferrosilicon alloy to give its special characteristic. The
background for development of this product has been based on new fundamental understanding of graphite nucleation
mechanisms in ductile iron, where the main body of nucleation sites was found to be comprised of complex but very stable
sulphides and oxides (Skaland 1992). Figure 1 shows an example of such nucleation site in ductile iron both as a high
magnification micrograph and a schematic representation of its phase composition. In conventional ductile iron production
the availability of such sulphide and oxide nucleation sites are determined by the purity of base metal and its additives,
holding times and temperature as well as metallurgical treatment processes and additives.
Traditionally, commercial inoculants have been based on ferrosilicon alloys containing metallic additives such as Calcium,
Barium, Strontium, Aluminum, Zirconium, Rare Earth’s, etc., with the main objective of these reactive elements to combine
with Sulphur and Oxygen in the iron and form potent heterogeneous nucleation sites for graphite. However, with restricted
availability of Sulphur and Oxygen in the iron, the metallic inoculant additives may reach a performance limit where their
effectiveness are restricted by the number of potent nucleation sites that can be formed after treatment. Thus, the primary
objective of the new inoculant concept has been to introduce controlled concentrations of non-metallic elements such as
Sulphur and Oxygen with the metallic inoculant. From balanced and controlled inclusion engineering, this will deliberately
produce a higher number density of nucleation sites for graphite from a reaction taking place between the highly reactive
metallic ingredients (Ca and Ce) and the non-metallic ingredients (S and O) of the inoculant. These additional nucleation
sites will then perform in parallel to the traditional sites formed during reactions between nodulizer, inoculant and the base
metal. The outcome will be a remarkable improvement in conditions for controlled graphite precipitation and growth, with all
possible benefits this may introduce to the final iron quality.
Several researchers have proven the benefits of Sulphur to graphite nucleation (Chisamera 1994, Lalich 1976, Mercier 1969).
Also, it has been proposed that Oxygen may play a vital role in the inoculation process (Tartera 1980, Nakae 1992, Podrzucki
2000). However, the combined use and performance of both elements through post-inoculation is a novel approach that has
been designed to get the benefits from both Calcium, Cerium, Sulphur and Oxygen simultaneously in the graphite nucleation
process. Calcium is used as the primary reactive element in inoculation, and has proven crucial for eutectic graphite
1
The new inoculant is available under the trade name Ultraseed® inoculant, produced by Elkem.
1
nucleation (Bilek 1972). Cerium is introduced for several reasons. First, Cerium will contribute in neutralizing subversive
trace elements in the base iron, forming stable inter-metallic compounds (Park 2000, Udomon 1985). Cerium will also have
strong affinity to Sulphur and Oxygen, resulting in the formation of highly stable Cerium oxides, sulphide, and oxy-sulphides
(Kozlov 1991, Warrick 1966). These Cerium compounds appear to be very beneficial in the inoculation process, resulting in
improved nucleation effectiveness throughout the entire solidification range.
(a)
(b)
Figure 1. (a) Example of duplex sulphide/oxide nuclei particle in ductile iron at large magnification in a
transmission electron microscope (70,000X). (b) Schematic representation of a nucleus particle containing
complex sulphide and oxide phases after nodularizing and inoculation of the iron (Skaland 1992).
EFFECTS OF INOCULATION ON CAST IRON PROPERTIES
The principal effects of cast iron inoculation can be described as follows:

Avoid the formation of hard carbides (cementite)
Promote the formation of graphite and ferrite
Reduce the segregation tendency of alloying and trace elements
Reduce the solidification shrinkage tendency
Improve the machinability of castings
Reduce the hardness
Increase the ductility
Give more homogeneous structures and properties in different sections of complex castings
The new inoculant concept is found to improve most all of these properties to a greater extent than other ferrosilicon based
inoculant alloys. Especially, improvements in ferrite formation, shrinkage minimizing, machinability, and microstructure
homogeneity have been observed through extensive testing in various foundry conditions. In the following, some of the
unique features will be described in more detail, also including examples from foundries. A series of case studies will be
reviewed to illustrate the performance characteristics through realistic examples from the industry.
UNIQUE FEATURES OF THE NEW INOCULANT CONCEPT
The new inoculant concept provides formation of extra nucleation sites in ductile iron in addition to those initially generated
by the magnesium treatment. This will increase nodule count and improve nodularity thus reducing carbide and shrinkage
tendency. The balanced cerium content also neutralizes subversive elements that may prevent the formation of nodular
graphite. Due to the higher nodule count obtained, the inoculant also provides formation of more ferrite in ductile irons. This
can be an advantage when producing the higher ductility and impact resistant grades of ferritic iron (e.g. grade 40.3).
The powerful nucleation characteristics are based on the formation of special cerium-calcium-sulphides and oxides that will
act as effective nucleation sites for graphite during solidification of the iron. These nucleation sites will contribute together
with the primary magnesium-silicon oxides to give powerful graphite nucleation with the outcome being a very high nodule
number density. The inoculant is found to be especially potent in ductile irons of relatively low sulphur content and in irons
treated with magnesium metal in a converter or wire injection process. The introduction of Cerium (Rare Earth metal)
through the inoculant can also replace the need for Rare Earth’s to be added through the nodularizing process. The inoculant
has also proven highly successful in providing fresh nucleation sites to ductile irons of long holding time where the base iron
2
or magnesium treated iron have been held for prolonged times before addition of the post inoculant. Such long hold times are
well known to reduce the overall nucleation capabilities of the iron prior to inoculation resulting in so-called “dead” iron. The
new inoculant concept will thus re-install good nucleation effectiveness from reactions with its deliberate Sulphur and
Oxygen content forming additional, new nucleation sites.
Due to the powerful effects on raising nodule count and improving chill protection, it has been found that the tendency to
shrinkage cavity formation is also greatly reduced with this inoculant. Especially, the type of shrinkage that often occur as
small porosities in hot-spot sections of complex castings appear to be effectively reduced or even eliminated by this inoculant
concept. It has been found that a characteristic bi-modal size distribution of nodules often will occur from a secondary, late
precipitation event in the last part of the solidification sequence. Such late graphite expansion effects will effectively
counteract shrinkage contraction in the last part of solidification, when risers have stopped functioning and graphite
expansion is most needed to counteract shrink. It appears that the new inoculant concept is effectively distributing the
graphite nucleation and growth phenomena throughout the entire solidification range. Conventional inoculants, however,
have a tendency to give massive and early expansion effects and very little contribution in the last part of solidification when
really most needed.
Strong nucleation effect and high nodule count is also a prerequisite to maximize the ferrite content when producing as-cast
ferritic grades of ductile iron. Particularly when there are limitations on the final silicon content of the iron, the high nodule
count obtained with this inoculant has proven effective to ensure the required minimum content of ferrite in such castings.
The bi-modal size distribution of nodules, and the fact that the smaller and late precipitated nodules are formed in the last
liquid to freeze, also aid in formation of more ferrite by acting as effective carbon sinks in these segregated areas enriched in
pearlite promoting elements. This is also indirectly improving the machinability of ductile iron, and the inoculant should
therefore be the preferred choice when good machinability is an important requirement.
The special inoculant composition including additions of finely dispersed oxides and sulphides with the ferrosilicon based
alloy, causes a specific appearance of this product. Figure 2 shows a comparison of the new inoculant design and a
conventional ferrosilicon based alloy. It is clear that the new concept appears with characteristic black particle surfaces due to
its sulphide and oxide content. Table 1 gives the specifications and typical composition of the new inoculant concept.
(a)
(b)
Figure 2. Physical appearance of (a) conventional ferrosilicon inoculant with metallic, glinsing surface
characteristics, (b) new inoculant concept with black surface characteristics.
Table 1. Specifications and typical composition of the new inoculant concept.
% Silicon
Specifications
Typicals
70 - 76
73
% Calcium
0.75 – 1.25
1.0
% Cerium
1.5 – 2.0
1.75
% Aluminum
0.75 – 1.25
1.0
% Sulphur
Max. 1.0
Trace
% Oxygen
Max. 1.0
Trace
3
RESULTS FROM FOUNDRY TESTING
The new inoculant concept has now been tested and implemented in numerous foundries Globally. More than 80 foundries
have conducted testing so far, and of these above 60% have reported some kind of successful results. Foundries have
different criteria and objectives in testing, and are looking for individual types of property improvements. The following
specific features may be mentioned from foundry testing:

Especially effective to reinstall powerful inoculation conditions in “dead” base irons
Improved performance both in pure Magnesium and Magnesium ferrosilicon treatment applications
Especially effective used as a late in-stream inoculation
Especially effective in eliminating shrinkage porosity in complicated hot-spot sections
Reduces the section sensitivity of nodule structures in castings of variable thickness
May not be that efficient as an all-round gray iron inoculant
Successful applications also observed in compacted graphite iron (reduced section sensitivity)
In the following, four case studies from foundry testing will be reviewed. The intention with these case studies is to show
some typical examples of performance characteristics observed in different foundry conditions with the new inoculant
concept.
CASE STUDY 1
This foundry uses electric induction melting and a tundish ladle process for preparing ductile iron. The treated ductile iron is
transferred into a fairly large channel induction holding and pouring unit where iron may sit for a while. The foundry has
experienced problems with carbides and excess shrinkage in complex castings of very thin sections. A key challenge is to
avoid big shrinkage porosity in a distant hot spot knob section that has to be drilled and need a smooth inner hole surface.
The foundry normally uses an in-stream late post inoculation addition of a (Zr,Mn,Ca)-bearing ferrosilicon inoculant.
When first testing the new inoculant concept, this was directly compared to the (Zr,Mn,Ca)-bearing ferrosilicon alloy as an
in-stream addition. Table 2 shows the typical average nodule number densities obtained for the two inoculants, while Figures
3 through 5 shows examples of the effects on nodule structure, carbide formation, and hot-spot shrinkage formation tendency.
It is evident from the results that a pronounced difference in nodule number density and size distribution occur for the two
inoculants. The new inoculant concept gives about double the nodule number density of the (Zr,Mn,Ca)-bearing inoculant,
and the nodule size distribution shows a shift to numerous smaller and better shaped nodules.
Figure 4 also shows the effects on carbide formation tendency in a very thin, 2-3 mm section of the same casting. The
(Zr,Mn,Ca)-bearing inoculant with its lower nodule count reveals the appearance of intercellular carbides, while the
(Ca,Ce,S,O)-inoculant has effectively eliminated these thin section carbides. Finally, Figure 5 shows the effects of
inoculation on the critical and difficult hot-spot section. The large region of micro-shrinkage occurring with the (Zr,Mn,Ca)-
bearing inoculant has been effectively minimized with the new (Ca,Ce,S,O)-inoculant. The extension of micro-porosity in
Figure 5b has proven small enough avoiding rough inner surfaces after drilling of this critical knob section.
Table 2. Average nodule number density for the (Zr,Mn,Ca)-bearing ferrosilicon inoculant
and the new inoculant concept.
(Zr,Mn,Ca)-inoculant
315 nodules per mm2
New inoculant concept
602 nodules per mm2
Conclusively, it can be said that this foundry is very happy with the new inoculant performances. A complete transition from
the (Zr,Mn,Ca)-bearing inoculant to the (Ca,Ce,S,O)-inoculant has been taking place for ductile iron production. Significant
improvements in reject rate and final casting quality has been obtained. The present addition rate of the new inoculant
concept is about 25% less than the previous (Zr,Mn,Ca)-bearing alloy, still providing the significant improvements.
4
(a)
(b)
(c)
(d)
Figure 3. Examples of microstructure results from Test Foundry 1.
(Zr,Mn,Ca)-containing inoculant, (a) polished condition, (b) etched in Nital.
New (Ca,Ce,S,O)-containing inoculant, (c) polished condition, (d) etched in Nital. (100X)
(a)
(b)
Figure 4. Examples of carbide conditions in thin 2 mm flange from Test Foundry 1.
(Zr,Mn,Ca)-containing inoculant revealing intercellular carbides.
(Ca,Ce,S,O)-containing inoculant showing carbide free conditions.
(a)
(b)
Figure 5. Examples of hot-spot micro-shrinkage porosity formation tendency from Test Foundry 1.
(a) (Zr,Mn,Ca)-containing inoculant causing massive shrinkage porosities.
(b) (Ca,Ce,S,O)-containing inoculant giving only traces of micro-shrinkage porosity.
5
CASE STUDY 2
This foundry is also an induction melting, sandwich treatment operation, and here the objective has been to test a series of
different generic inoculants in order to find the optimum product for their autopouring and in-stream inoculation on a DISA
molding line. Output parameters evaluated include nodule- and microstructure, mechanical properties and shrinkage tendency
for the different inoculants.
(a)
(b)
(c)
(d)
Figure 6. Examples of graphite nodule structure in plate castings from Test Foundry 2.
5 mm section size: (a) Sr-containing inoculant, (b) (Ca,Ce,S,O)-containing inoculant.
40 mm section size: (c) Sr-containing inoculant, (d) (Ca,Ce,S,O)-containing inoculant (100X).
This foundry normally uses a Strontium-bearing ferrosilicon inoculant, and Figure 6 shows an example of the effects of Sr-
FeSi and (Ca,Ce,S,O)-FeSi inoculation on the final nodule structure in a thin 5 mm and a thick 40 mm plate section casting.
A quite remarkable difference is observed, especially for the heavier 40 mm plate section, where the (Ca,Ce,S,O)-inoculant
appears to give a very strong increase in nodule count. Figure 7 shows a quantitative comparison of nodule counts for the two
inoculants in sections of 5, 10, 20 and 40 mm thickness. The Sr-FeSi inoculant gives a quite normal and expected behavior
with a falling nodule count from about 300 per mm2 to 150 per mm2 when the section thickness increases from 5 through 40
mm. This behavior is normally observed for most commercial inoculants when increasing the section thickness.
The (Ca,Ce,S,O)-containing inoculant on the other hand, shows a complete different and quite remarkable behavior. The
histograms in Figure 7 and also the micrographs in Figure 6, shows an almost unaffected nodule count for the spread in
section thickness. About 300 nodules per mm2 are measured for both the 5 and 40 mm sections of the same test casting. In
fact, the 40 mm section contains an even higher nodule number density of 340 per mm2 versus only 312 per mm2 for the
much thinner 5 mm section. This is clearly confirmed by the micrographs in Figures 6b and 6d.
6
400
350
300
250
N/mm2
200
150
100
50
0
5
10
20
40
Ce, S, O
Sr
Section size (mm)
Figure 7. Example of nodule count in various section thicknesses from 5 to 40 mm plate castings from
Test Foundry 2 for (Ca,Ce,S,O)-containing inoculant and Sr-containing inoculant.
This unusual behavior opens up for some interesting performances. First, it is evident that the section sensitivity for complex
castings can be greatly reduced and microstructure and properties controlled and equalized for a large span in section
thickness. This in itself may offer advantages to the homogeneous production of difficult castings with demanding properties
in different locations. Further, the observation of an elevated nodule count in heavier sections also offer some additional
benefits in relation to solidification contraction and shrinkage tendency.
Figure 8 shows examples of crossbar hot spot shrinkage conditions using three different inoculants in Test Foundry 2. The
inoculants are Ba-containing, Sr-containing, and the new (Ca,Ce,S,O)-containing FeSi. As the Figure shows, shrinkage
tendency differs greatly for the different inoculants. Both the Ba-containing and Sr-containing alloys give massive
contraction effects and large cavities in the hot-spot cross bar. The (Ca,Ce,S,O)-containing inoculant on the other hand,
shows an almost complete elimination of shrinkage porosity with only one very small cavity revealed in the section cut
through the parting line of the experimental cross bar casting. This dramatic effect on shrinkage tendency can be directly
related to the nodule formation and the rate of graphite growth throughout the entire solidification sequence.
As shown in Figures 6 and 7, the conventional inoculant behavior, represented by the Sr-inoculant, is to give fairly uniform
nodule sizes and a reduction in nodule count as the section size increases. With the (Ca,Ce,S,O)-inoculant concept, there is an
effect causing a bi-modal nodule size distribution and numerous smaller nodules that are precipitated very late during
solidification. This late graphite expansion effectively counteract shrinkage contraction, as can be clearly seen in Figure 8c
for the (Ca,Ce,S,O)-containing inoculant. Both the Sr- and Ba-containing inoculants give low nodule counts for heavier
sections, around 200-220 per mm2 , while the (Ca,Ce,S,O)-containing inoculant gives about 350 nodules per mm2 for the
similar section.
Since this phenomenon predominantly occur for heavier sections, this is where the bi-modal size distribution is most clearly
observed and also where shrinkage control is mostly needed. There exist no clear understanding of the mechanisms of the late
graphite formation and the resulting bi-modal nodule distribution effect. However, it is expected that the introduction of
Cerium in combination with sulphur and Oxygen in the inoculant, will introduce more nucleation sites, and possibly also a
second type of sites that are activated later during solidification. Cerium oxides and oxy-sulphides will behave very different
from the traditional Ca,Ba,Sr-type oxides and silicates know to nucleate primary graphite in the early stages of solidification
(see Figure 1b).
7
(a)
(b)
(c)
Figure 8. Examples of shrinkage porosity formation in crossbar castings from Test Foundry 2.
(a) Ba-containing inoculant, (b) Sr-containing inoculant, (c) (Ca,Ce,S,O)-containing inoculant.
Since the phenomenon is most evident in heavier sections, it is likely that this second type of beneficial and late activated
nucleation sites only show their characteristic effects in slower cooling conditions. This is, the second type nucleation sites
need more time to become activated, and will thus only give maximum benefits in heavier sections of a casting. The net
outcome appears as a uniform nodule count in different sections, combined with effective shrink elimination in the heavier
and slower cooled sections.
The conclusion from this extensive inoculant testing has been that the Test Foundry 2 now has converted to use the new
(Ca,Ce,S,O)-containing inoculant for all their ductile iron production. Great improvements in especially shrinkage reduction
have been experienced, and the inoculant addition rate also reduced to a minimum.
CASE STUDY 3
The third case study represents a foundry producing very heavy ductile iron castings. Induction melting and tundish treatment
is applied also here. In this case, the foundry suffers from the classical problems of heavy section castings such as graphite
flotation, segregation, shrinkage and relatively poor nodularity. The foundry uses manual transfer inoculation to large pouring
ladles, and the present inoculant material applied is a Barium-containing ferrosilicon alloy. Barium inoculants are
traditionally recommended for heavy casting and slow cooling applications, since Ba is recognized for its minimum of fading
tendency during prolonged hold and solidification times. The new (Ca,Ce,S,O)-type inoculant was tested in parallel to the
Ba-inoculant as a ladle addition, and effects on microstructure, machinability and shrinkage tendency evaluated.
Figure 9 shows examples of typical microstructures obtained with the two different inoculants in a fairly heavy 50 mm
section. The Ba-inoculant shows the expected relatively large and uniformly sized nodules in a ferritic/pearlitic matrix (see
Figures 9a and 9b). The (Ca,Ce,S,O)-inoculant, on the other hand, shows a much wider spread in nodule sizes, and the
characteristic bi-modal distribution effect is again clearly revealed. Figures 9c and 9d shows the effect of the (Ca,Ce,S,O)-
inoculant on nodule distribution and ferrite/pearlite ratio. Table 3 also gives the quantified microstructure data for the two
respective inoculants in this heavy section application.
From Figure 9 and Table 3 it is evident that the heavy section impact on microstructure for the (Ca,Ce,S,O)-inoculant is
significant. The bi-modal nodule distribution effect was found to effectively minimize difficult and massive shrinkage
formation in large castings. The formation of smaller nodules also gave a general improvement of 10% in the nodularity from
about 80 to 90 %. Further, Figures 9b versus 9d clearly show a significant reduction in intercellular pearlite with the new
inoculant concept. The reduction is quantified from 25% down to 13% pearlite. The interconnected network of pearlite at
25% and higher is broken down into only minor fragments of pearlite in a predominantly ferritic matrix. This is again due to
the numerous smaller nodules arising in the segregated intercellular regions, acting as carbon sinks during the eutectoid
transformation. Raising nodule counts by a general increase in the primary formed larger nodules, typically will not influence
the pearlite ratio to the same great extent. This is because segregation patterns and profiles will still remain the same. When
the grain boundary nodules are included, this will have a pronounced effect on scavenging the matrix for carbon, thus
effectively reducing the risk for harmful segregation phenomena and formation of intercellular carbides, phosphides, and
other unwanted microconstituents. A general improvement in tensile and impact properties was also experienced with the bi-
modal and homogeneous nodule distribution. Also, an improved tool life during machining of up to 50% was experienced
with this new situation.
8
(a)
(b)
(c)
(d)
Figure 9. Examples of microstructure in heavy section castings from Test Foundry 3.
Ba-containing inoculant:(a) polished condition, (b) etched in Nital.
(Ca,Ce,S,O)-containing inoculant:(c) polished condition, (d) etched in Nital.
Table 3. Effects of Ba- and (Ca,Ce,S,O)-containing inoculants on microstructure characteristics in
heavy section casting at Test Foundry 3.
Nodule count
Per mm2
Ba-inoculant
(Ca,Ce,S,O)-inoculant
187
357
Nodularity
%
80
90
Pearlite
%
25
13
Shrinkage
tendency
Significant
Much less
Relative
machinability
Medium
Good
CASE STUDY 4
The final case study included here is from an automotive foundry using electric induction melting and sandwich magnesium
treatment. Molding is done on BMD and DISA lines, autopouring through Junker units. Late in-stream inoculation is applied
on all pouring lines, and typically Zirconium-bearing inoculants have been applied. The foundry is suffering from some
serious shrinkage problems, and the type shrink can be described as “massive” cavities in critical sections. Extensive risering
has been implemented to try and overcome the problems, but still large shrinkage cavities are found to occur even adjacent to
the risers.
The new inoculant concept using a (Ca,Ce,S,O)-bearing ferrosilicon inoculants was tested out on the DISA line for a special
test pattern involving a square cubic test piece attached to a fairly large riser. The cube and riser are cut in half, and evaluated
for degree of shrinkage porosity formation and distribution of cavities in test casting and riser.
Figures 10 and 11 shows examples of such test castings cut through the middle for evaluation of shrink. The examples
compare two different Zirconium-containing inoculants, one (Zr,Mn,Ca)-type and one (Zr,Ca)-type, to the new (Ca,Ce,S,O)-
9
concept inoculant. Figure 10 shows conditions for the three inoculants, where the (Zr,Mn,Ca)-type to the left reveals a large
cluster of micro-shrinkage porosity in the body of the cube test casting. The riser is virtually sound. The (Zr,Ca)-inoculant in
the middle shows one large cavity in the cube sample and another in the neck of the riser, while the (Ca,Ce,S,O)-inoculant to
the right shows only a very small porosity in the cube test piece.
Conditions shown in Figure 11 compares the (Zr,Ca)-inoculant to the (Ca,Ce,S,O)-inoculant. The test piece for the (Zr,Ca)-
inoculant to the left shows again massive shrinkage in the cube sample and still quite widespread porosities also in the riser.
The (Ca,Ce,S,O)-inoculant to the right has pushed the shrinkage void all the way back to the top of the riser, leaving the
lower part of the riser and the cube casting itself sound and completely free from porosities.
The foundry has solved a severe shrinkage problem using the new concept inoculant. Complete conversion from Zr-bearing
to (Ca,Ce,S,O)-bearing inoculation has been implemented at this foundry.
Figure 10. Example of cubic test castings with attached riser from Test Foundry 4.
Inoculants tested are: Left: (Zr,Mn,Ca)-bearing, middle: (Zr,Ca)-bearing, and
right: (Ca,Ce,S,O)-bearing.
There exist numerous other cases showing similar improvements in chill situation, nodule structure, and shrinkage formation
tendency using the new concept (Ca,Ce,S,O)-containing inoculant. However, from space limitations this paper is restricted to
only cover a selected handful of classical situations experienced in small/thin and heavy/thick casting conditions. As shown
above, effects are found to be most pronounced for nodule count and size distribution, as well as carbide restriction and
shrink control. Other foundries have also reported great improvements in machinability conditions, and one foundry even
eliminated their heat treatment operation after converting inoculants. Double tool life at half the addition rate of inoculant
versus the previous regular calcium-bearing ferrosilicon has also been reported. Some foundries report of increased tensile
strength even with a significant increase in the ferrite content.
This paper most of all shows that the choice of inoculant material is not a trivial thing, and that different commercial
inoculants may have dramatic effects on the final ductile iron quality. It is important to keep in mind, that sometimes the least
expected inoculant is the one to actually perform best. Systematic and thoroughly controlled foundry testing will be the only
sound and safe way to ensure that the optimum cost efficient alternative inoculant is being used in the individual foundry.
There are too many unknown and uncontrollable factors affecting inoculation to give a general recommendation, and testing
to solve specific challenges will at the end be the only safe way to find the improved or optimized inoculant solution.
10
Figure 11. Example of cubic test castings from foundry 4. Inoculants tested are:
Left: (Zr,Ca)-bearing, right: (Ca,Ce,S,O)-bearing. Note the differences in shrinkage
cavity distribution for the two inoculants.
When testing inoculants in cast iron it is always important not only to look for a quick micrograph. Very attractive effects
may then be overseen. A thorough evaluation of all the above discussed factors must be considered, since great savings may
be found in reduced scrap rate from shrink, poor structures, machinability or even tensile and impact properties.
CONCLUSIONS
The following main conclusions can be given from the present investigation:

A new approach to ductile iron inoculant design has been described. The new design has proven successful in improving
casting performance and properties. The new ferrosilicon based inoculant material contains levels of Calcium and
Cerium that are adjusted to minimize chill formation and neutralize subversive trace elements in the iron.
The new inoculant design also contains small and controlled amounts of Sulphur and Oxygen in a form that make them
available for reaction with the Calcium and Cerium during introduction into liquid iron. The special composition is
designed to give highly powerful graphite nucleation conditions in ductile irons along with very effective chill and
shrinkage reduction.
Experience from foundry testing has proven that the new inoculant concept is especially effective in re-installing
powerful inoculation conditions in irons of a “dead” nature. Also, especial effectiveness in minimizing shrinkage
porosity in complicated hot-spot sections has been observed.
The new (Ca,Ce,S,O)-containing inoculant is also reducing the section sensitivity of nodule structures in castings of
variable thickness. A bi-modal size distribution of graphite nodules is often observed. This nodule distribution is
instrumental in minimizing shrink and intercellular pearlite and carbide formation.

11
REFERENCES
Bilek, P.J., Dong, J.M., McCluhan, T.K., “The Role of Ca and Al in Inoculation of Gray Iron”, AFS Transactions, pp.183-
188, (1972)
Chisamera, M., Riposan, I., Proc. 5th Int. Symp. On the Physical Metallurgy of Cast Iron, Nancy, France, Sept. (1994)
Kozlov, L.J., Vorobyev, A.P., “The Role of Rare-earth Metals in the Process of Spheroidal Graphite Formation”, Cast
Metals, vol.4, no.1, (1991)
Lalich, M.J., Hitchings, J.R., “Characterization of Inclusions as Nuclei for Spheroidal Graphite in Ductile Cast Iron”, AFS
Transactions, pp.653-664, (1976)
Mercier, J-C., Paton, R., Margerie, J-C., Mascre, C., “Inclusions dans les spheroides de graphite”, Fonderie, April (1969)
Nakae, H., Koizumi, H., Takai, K., Okauchi, K., “Nucleation of Graphite in Inoculated Cast Iron”, Trans. Japan
Foundrymen’s Society, vol.11, pp. 34-39, (1992)
Park, J., Loper, C.R., “Neutralizing of Lead in Gray Iron Melts Using Misch Metal”, AFS Transactions, (2000)
Podrzucki, C., Fras, E, Lopez, H.F., “The Inoculation of Cast Iron: Role of Oxygen”, AFS Transactions, (2000).
Skaland, T., A Model for the Graphite Formation in Ductile Cast Iron, Ph.D. Thesis 1992:33, The University of Trondheim,
NTH, Department of Metallurgy, Norway, (1992)
Tartera, J., “Cast Iron Inoculation Mechanisms”, AFS Int. Cast Metals J., pp.7-14, December (1980)
Udomon, U.H., Loper, C.R., “Comments Concerning the Interaction of Rare Earths With Subversive Elements In Cast
Irons”, AFS Transactions, pp.519-522, (1985)
Warrick, R.J., “Spheroidal Graphite Nuclei in Rare Earth and Magnesium Inoculated Irons”, AFS Cast Metals Research J.,
pp.97-108, Sept., (1966)
12

A NEW METHOD FOR CHILL AND SHRINKAGE CONTROL

2009-07-21T21:29:00.000-07:00

در پيام بعدي مقاله
A NEW METHOD FOR CHILL AND SHRINKAGE CONTROL
IN LADLE TREATED DUCTILE IRON
را خدمتتان معرفي خواهم كرد.هر مشكل متالورژيكي كه داشته با شيد با ما در جريان قرار دهيد

A New Approach to Ductile Iron Inoculation

2009-07-21T21:24:00.000-07:00

A New Approach to Ductile Iron Inoculation

T. Skaland
Elkem ASA, Research
Kristiansand, Norway
Copyright©2001 American Foundry Society
ABSTRACT
The objective of the present paper is to describe a new approach to inoculant design that has proven successful in improving
casting performance and properties. The described inoculant represents a unique new generation of products developed for
powerful cast iron inoculation. The ferrosilicon alloy contains levels of Calcium and Cerium that are adjusted to minimize
chill formation and neutralize subversive trace elements in the iron. In addition, the inoculant contains small and controlled
amounts of Sulphur and Oxygen in a form that make them available for reaction with the Calcium and Cerium during
introduction into liquid iron. This special composition is designed to give highly powerful graphite nucleation conditions in
ductile irons along with very effective chill and shrinkage reduction. Examples from foundry testing are reviewed, and the
unique characteristics of this new inoculant concept described. The new inoculant is found to be more powerful than
conventional ferrosilicon based inoculants, and give rise to very effective reduction in the shrinkage tendency of ductile irons.
Special effectiveness has been observed in irons of low sulphur or irons of a “dead” nature from prolonged holding times.
Also, results show improvements in both tensile properties as well as machinability for ferritic ductile irons. The new
inoculant concept represents a patent protected design (Int.Pat.No.WO99/29911), and is available under a special trade name,
see footnote1 .
INTRODUCTION
Based on years of laboratory work with experimental inoculants of various compositions, a new concept for inoculation of
ductile iron has been developed. The concept is novel in the sense that it involves not only an alloyed ferrosilicon-based
material, but also introduction of non-metallic powders with the ferrosilicon alloy to give its special characteristic. The
background for development of this product has been based on new fundamental understanding of graphite nucleation
mechanisms in ductile iron, where the main body of nucleation sites was found to be comprised of complex but very stable
sulphides and oxides (Skaland 1992). Figure 1 shows an example of such nucleation site in ductile iron both as a high
magnification micrograph and a schematic representation of its phase composition. In conventional ductile iron production
the availability of such sulphide and oxide nucleation sites are determined by the purity of base metal and its additives,
holding times and temperature as well as metallurgical treatment processes and additives.
Traditionally, commercial inoculants have been based on ferrosilicon alloys containing metallic additives such as Calcium,
Barium, Strontium, Aluminum, Zirconium, Rare Earth’s, etc., with the main objective of these reactive elements to combine
with Sulphur and Oxygen in the iron and form potent heterogeneous nucleation sites for graphite. However, with restricted
availability of Sulphur and Oxygen in the iron, the metallic inoculant additives may reach a performance limit where their
effectiveness are restricted by the number of potent nucleation sites that can be formed after treatment. Thus, the primary
objective of the new inoculant concept has been to introduce controlled concentrations of non-metallic elements such as
Sulphur and Oxygen with the metallic inoculant. From balanced and controlled inclusion engineering, this will deliberately
produce a higher number density of nucleation sites for graphite from a reaction taking place between the highly reactive
metallic ingredients (Ca and Ce) and the non-metallic ingredients (S and O) of the inoculant. These additional nucleation
sites will then perform in parallel to the traditional sites formed during reactions between nodulizer, inoculant and the base
metal. The outcome will be a remarkable improvement in conditions for controlled graphite precipitation and growth, with all
possible benefits this may introduce to the final iron quality.

بزرگتري وبلاگ خدمات متالورژي

2009-07-17T23:59:00.000-07:00

با سلام از اين به بعد در اين وبلاگ خدمات مورد نياز ريخته گران عزير اعم از مدل سازي ،قالب سازي،ريخته گري،اطلاع رساني و تامين منابع اطلاعاتي ،عمليات حرارتي و ...در سراسر كشور ارائه خواهد شد حتما با ما در تماس و ارتباط باشيد.