JP3775335B2 - Silicon nitride sintered body, method for producing silicon nitride sintered body, and circuit board using the same - Google Patents
Silicon nitride sintered body, method for producing silicon nitride sintered body, and circuit board using the same Download PDFInfo
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- JP3775335B2 JP3775335B2 JP2002121345A JP2002121345A JP3775335B2 JP 3775335 B2 JP3775335 B2 JP 3775335B2 JP 2002121345 A JP2002121345 A JP 2002121345A JP 2002121345 A JP2002121345 A JP 2002121345A JP 3775335 B2 JP3775335 B2 JP 3775335B2
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- silicon nitride
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- 229910052581 Si3N4 Inorganic materials 0.000 title claims description 189
- HQVNEWCFYHHQES-UHFFFAOYSA-N silicon nitride Chemical compound N12[Si]34N5[Si]62N3[Si]51N64 HQVNEWCFYHHQES-UHFFFAOYSA-N 0.000 title claims description 188
- 238000004519 manufacturing process Methods 0.000 title claims description 17
- 239000002245 particle Substances 0.000 claims description 99
- 238000005245 sintering Methods 0.000 claims description 57
- 239000000843 powder Substances 0.000 claims description 49
- 229910052760 oxygen Inorganic materials 0.000 claims description 33
- QVGXLLKOCUKJST-UHFFFAOYSA-N atomic oxygen Chemical compound [O] QVGXLLKOCUKJST-UHFFFAOYSA-N 0.000 claims description 32
- 239000001301 oxygen Substances 0.000 claims description 32
- 229910052761 rare earth metal Inorganic materials 0.000 claims description 30
- 239000011777 magnesium Substances 0.000 claims description 25
- 238000010438 heat treatment Methods 0.000 claims description 22
- CPLXHLVBOLITMK-UHFFFAOYSA-N magnesium oxide Inorganic materials [Mg]=O CPLXHLVBOLITMK-UHFFFAOYSA-N 0.000 claims description 21
- 239000000395 magnesium oxide Substances 0.000 claims description 21
- AXZKOIWUVFPNLO-UHFFFAOYSA-N magnesium;oxygen(2-) Chemical compound [O-2].[Mg+2] AXZKOIWUVFPNLO-UHFFFAOYSA-N 0.000 claims description 21
- 229910001404 rare earth metal oxide Inorganic materials 0.000 claims description 16
- IJGRMHOSHXDMSA-UHFFFAOYSA-N Atomic nitrogen Chemical compound N#N IJGRMHOSHXDMSA-UHFFFAOYSA-N 0.000 claims description 11
- 229910052757 nitrogen Inorganic materials 0.000 claims description 7
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- YXFVVABEGXRONW-UHFFFAOYSA-N Toluene Chemical compound CC1=CC=CC=C1 YXFVVABEGXRONW-UHFFFAOYSA-N 0.000 description 3
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- CMIHHWBVHJVIGI-UHFFFAOYSA-N gadolinium(III) oxide Inorganic materials [O-2].[O-2].[O-2].[Gd+3].[Gd+3] CMIHHWBVHJVIGI-UHFFFAOYSA-N 0.000 description 3
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- LIVNPJMFVYWSIS-UHFFFAOYSA-N silicon monoxide Chemical class [Si-]#[O+] LIVNPJMFVYWSIS-UHFFFAOYSA-N 0.000 description 3
- NIQCNGHVCWTJSM-UHFFFAOYSA-N Dimethyl phthalate Chemical compound COC(=O)C1=CC=CC=C1C(=O)OC NIQCNGHVCWTJSM-UHFFFAOYSA-N 0.000 description 2
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- LRHPLDYGYMQRHN-UHFFFAOYSA-N N-Butanol Chemical compound CCCCO LRHPLDYGYMQRHN-UHFFFAOYSA-N 0.000 description 2
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Description
【０００１】
【発明の属する技術分野】
本発明は、半導体用基板や発熱素子用ヒ−トシンク等の電子部品用部材、あるいは一般機械器具用部材、溶融金属用部材、または熱機関用部材等の構造用部材として好適な高強度・高熱伝導性に富んだ窒化ケイ素質焼結体およびその製造方法、ならびに前記窒化ケイ素質焼結体を用いて構成される回路基板に関する。
【０００２】
【従来の技術】
窒化ケイ素質焼結体は、高温強度特性および耐摩耗性等の機械的特性に加え、耐熱性、低熱膨張性、耐熱衝撃性、および金属に対する耐食性に優れているので、従来からガスタ−ビン用部材、エンジン用部材、製鋼用機械部材、あるいは溶融金属の耐溶部材等の各種構造用部材に用いられている。また、高い絶縁性を利用して電気絶縁材料として使用されている。
【０００３】
近年、高周波トランジスタ、パワーＩＣ等の発熱量の大きい半導体素子の発展に伴い、電気絶縁性に加えて良好な放熱特性を得るために高い熱伝導率を有するセラミックス基板の需要が増加している。このようなセラミックス基板として、窒化アルミニウム基板が用いられているが、機械的強度や破壊靭性等が低く、基板ユニットの組立て工程での締め付けによって割れを生じるという問題がある。また、Ｓｉ半導体素子を窒化アルミニウム基板に実装した回路基板では、Ｓｉと窒化アルミニウム基板との熱膨張差が大きいため、熱サイクルにより窒化アルミニウム基板にクラックや割れを発生し実装信頼性が低下するという問題がある。
【０００４】
そこで、窒化アルミニウム基板より熱伝導率は劣るものの、熱膨張率がＳｉに近く、かつ機械的強度、破壊靭性および耐熱疲労特性に優れる高熱伝導窒化ケイ素質焼結体からなる基板が注目され、種々の提案が行われている。
【０００５】
例えば、特開平４−１７５２６８号公報には、実質的に窒化ケイ素からなり、不純物として含有されるAlおよび酸素が共に3.5重量％以下であり、密度が3.15Mg/m３（3.15g/cm３）以上であり、40W/(m・K) 以上の熱伝導率を有する窒化ケイ素質焼結体が記載されている。
【０００６】
また、特開平９−３０８６６号公報には、85〜99重量％のβ型窒化ケイ素粒と残部が酸化物または酸窒化物の粒界相とから構成され、粒界相中にMg,Ca，Sr，Ba，Y，La，Ce，Pr，Nd，Sm，Gd，Dy，Ho，ErおよびYbのうちから選ばれる少なくとも１種の元素を0.5〜10重量％含有し、粒界相中のAl元素含有量が1重量％以下であり、気孔率が5％以下であり、かつβ型窒化ケイ素粒のうちで短軸径5μｍ以上を持つものの割合が10〜60体積％である窒化ケイ素質焼結体が記載されている。
【０００７】
また、日本セラミックス協会1996年年会講演予稿集1G11、同1G12、および特開平１０−１９４８４２号公報には、原料粉末に柱状の窒化ケイ素粒子またはウイスカーを予め添加し、ドクターブレード法あるいは押出成形法を用いて、この粒子を２次元的に配向させた成形体を形成し、焼結することにより熱伝導に異方性を付与して特定方向の熱伝導率を高めた窒化ケイ素質焼結体が記載されている。
【０００８】
また、特開２００１−１９５５７号公報には、窒化ケイ素質粒内の酸素、Al、Ca、Fe、の不純物量の合計を1500ppm以下、かつ短軸径が2μｍ以上に制御することで、熱伝導率と機械特性を向上させた窒化ケイ素質焼結体ならびにその製造方法が記載されている。
【０００９】
さらに、特開２００２−２９８４８号公報には、原料粉末に柱状のウイスカーを予め添加し、焼成過程において当該ウイスカーを核として選択的に粒成長させたミクロ組織を構築することで、熱伝導率を向上させた窒化ケイ素質焼結体ならびに当該ウイスカーの製造方法が記載されている。
【００１０】
日本セラミックス協会1998年年会講演予稿集2B04には、窒化ケイ素粉末の成形体を1.0MPaの窒素ガス中で2000℃×4hrで焼結した後に、さらに30MPaの窒素ガス中で2200℃×4hrの高温高圧での熱処理を行うことにより、100w/(m・K)以上の高い熱伝導率を有する窒化ケイ素質焼結体が製造できることを記載している。これには、高熱伝導化の発現は焼結体中の窒化ケイ素粒子の成長に加えて、高温熱処理による窒化ケイ素粒子内での六角形の析出相が関与していると記載されている。
【００１１】
【発明が解決しようとする課題】
以下、上記した従来技術の問題点について順を追って説明する。
まず、特開平４−１７５２６８号公報では40W/(m・K)以上の熱伝導率が得られているが、昨今ではさらに熱伝導率を高めた、機械的強度に優れる材料が望まれている。
また、特開平９−３０８６６号公報、特開平１０−１９４８４２号等公報に記載の方法では、窒化ケイ素質焼結体中に巨大な柱状粒子を得るために、成長核となる種結晶あるいはウィスカ−を予め添加し、2000℃以上および10.1MPa（100気圧）以上の窒素雰囲気下での焼成が不可欠である。したがって、ホットプレスあるいはHIP等の特殊な高温・高圧設備が必要となりコストアップを招来する。また、窒化ケイ素粒子を配向させた成形体を得るための成形プロセスが複雑であるため、生産性が著しく低下するという問題がある。
【００１２】
また、特開２００１−１９５５７号公報に記載される窒化ケイ素質焼結体は、窒化ケイ素粒内の不純物量を低減する（純化効果と表記）ことで、粒子自身の熱伝導率を向上させ、これにより焼結体の熱伝導率を向上させることを特徴としている。また、この純化効果の助長に役立つ添加物として、Zrおよび／または、Hfを選定し、これらを酸化物換算として0.5wt%〜3.0wt%添加するとしている。しかしながら、ZrおよびHfの酸化物を添加すると、焼結過程でSi3N4中のNと容易に反応して粒界相中に電気伝導性のあるZrNおよびHfNが生成される。よって、本来、セラミックス基板に必須とされる電気絶縁性が保持できず、高周波で作動するパワー半導体モジュール用の絶縁基板として使用し難いという問題がある。
【００１３】
さらに、特開２００２−２９８４８号公報に記載される製造方法は、構成する窒化ケイ素粒子の平均円形度が、0.8以上あり、β化率が10％以上80％未満、酸素量が0.5〜1.8質量％、比表面積が12〜22ｍ2/ｇである窒化ケイ素粉末に、希土類酸化物、酸化珪素、および酸化マグネシウムよりなる群から選ればれる１種以上を、合計が2.5〜14質量％となるように添加し、更に、窒化珪素ウィスカ−を0.1〜8.5質量％添加した後、混合し、成形し、窒化雰囲気下で焼結させるものである。更に、焼結性向上のために窒化珪素ウィスカ−に対して、水沸点以上（例えば、110℃〜140℃の温度範囲下）で予め水熱処理することを特徴としている。しかしながら、この例で使用される窒化ケイ素ウィスカーの水熱処理は、表面を酸化させ助剤として作用するSiO2成分を増加させることで焼結性が改善できるが、その反面、実施例に記載がある様に緻密な焼結体を得るための水熱処理は、120℃で96ｈの処理が必要となりプロセスが煩雑になる。また、水熱処理により焼結性は改善できるものの、焼結体の酸素量、しいては窒化ケイ素粒子内の酸素量を低減することができず、高熱伝導材が得られ難いと言う難点がある。
【００１４】
次に、前述の日本セラミックス協会1998年年会講演予稿集2B04に記載の焼結体は、1MPa窒素ガス中2000℃での焼成後に、さらに30MPa窒素ガス中2200℃での高温高圧の熱処理を行うことにより100W/(m・K)以上の高い熱伝導率が得られる利点がある。更に、高熱伝導化の発現のメカニズムを、焼結体中の窒化ケイ素粒子の成長に加えて、高温熱処理によって窒化ケイ素粒子内の六角形の析出相が関与していると説明している。すなわち、焼結および粒成長時にY-Nd-Si-Oから構成される助剤成分が窒化ケイ素粒子内に取り込まれて固溶し、高温での熱処理および冷却時にY-Nd-Si-O組成のアモルファス相として、窒化ケイ素粒子内に析出し、析出物の一部は結晶化したものと考え、窒化ケイ素粒子の高純度化作用の１つとして考えられている。以上のことから上記の焼結体を得るには、特殊な高温・高圧設備が必要となりコストアップを招来する。更に焼結した上に熱処理を加えるため生産性が著しく低下するという問題がある。また、上記焼結体中の窒化ケイ素粒子内の析出相について詳細な組成分析ならびに観察がなされておらず、熱伝導率向上との関連性が明確にはなっていない。
【００１５】
本発明は上記従来の問題点に鑑みてなされたものであり、2000℃以上でかつ10.1MPa（100気圧）以上の高温・高圧焼成といったコストの高い焼成法を必要とせず、機械的強度に優れ、熱伝導の方向に異方性を持たずに従来に比べて熱伝導率を高めた高熱伝導型窒化ケイ素質焼結体を提供することを目的とする。
また本発明は、窒化ケイ素粒子内に析出する微細粒子の組成と形態を詳細に調査することにより熱伝導率を高めた高熱伝導型窒化ケイ素質焼結体を提供することを目的とする。
また本発明は、窒化ケイ素質粉末のβ分率、含有酸素量、不純物量およびα型窒化ケイ素質粉末との混合比及び保持過程を含む焼結工程等を規定することにより、高い熱伝導率と高い強度を有する窒化ケイ素質焼結体およびその製造方法を提供することを目的とする。
また本発明は、上記した高強度・高熱伝導性に富んだ窒化ケイ素質焼結体用いて構成される放熱性の良好な回路基板を提供することを目的とする。
【００１６】
【課題を解決するための手段】
本発明者らは上記課題を達成するため、窒化ケイ素粒子内に少なくとも酸素および焼結助剤成分を組成に含む微細粒子を意識的に析出させることで、窒化ケイ素粒子自身の熱伝導率を向上させ、安定して100W/(m・K)以上の熱伝導率と十分な曲げ強度を有する窒化ケイ素質焼結体が得られることを知見した。また、このとき焼結助剤成分はＭｇＯ基とすることで焼結性が向上し、かつＭｇＯと（ＲＥｘＯｙ）が特定量と特定比を持って含有していることが有効なことを知見した。また、上記窒化ケイ素質焼結体の製造方法においては、用いる窒化ケイ素質粉末のβ分率、含有酸素量、不純物およびα粉末との混合比等の粉末の特性及び保持過程を含む焼結工程等を規定することが肝要であることを知見した。以上により本発明に至ったものである。
【００１７】
即ち、本発明の窒化ケイ素質焼結体は、希土類元素（ＲＥ）からなる群から選ばれた少なくとも１種の元素とＭｇとが焼結助剤の成分として添加された窒化ケイ素質焼結体であって、前記希土類元素を希土類酸化物（ＲＥ ｘ Ｏ ｙ ）に換算し、前記Ｍｇを酸化マグネシウム（ＭｇＯ）に換算したとき、これら酸化物に換算した酸化物含有量の合計が 0.6 〜 10 ｗｔ％で、かつ（ＲＥ ｘ Ｏ ｙ ）／（ＭｇＯ）で表される重量比が１より大きく、前記焼結助剤の成分、酸素、窒素、及びシリコンが含まれる、平均粒径 100nm 以下の微細粒子が前記窒化ケイ素質焼結体を構成する窒化ケイ素粒子内に析出していることを特徴とする。また、本発明の窒化ケイ素質焼結体は、倍率 10,000 倍以上の透過型電子顕微鏡写真において、前記窒化ケイ素粒子内に平均粒径100nm以下の微細粒子が５個 / μ m 2 以上析出していることを特徴としている。前記微細粒子は、シリコン成分の高い核と、前記焼結助剤成分の高い周辺部とからなることを特徴とする。また、本発明の窒化ケイ素質焼結体は、前記微細粒子が非晶質相であることを特徴とする。当該微細粒子は、焼成過程で窒化ケイ素粒子の粒成長とともに極微量ではあるが粒内に取り込まれた助剤成分が、窒化ケイ素粒子内に再析出したものであり、窒化ケイ素粒子自身の高熱伝導化に寄与する。このとき、透過型電子顕微鏡（ＴＥＭ）による直接倍率10,000倍以上の観察像において、窒化ケイ素粒子内に平均粒径100nm以下の前記微細粒子が５個/μm2以上存在することが望ましく、この微細粒子の析出現象と割合により焼結体の熱伝導率は向上する。
【００１８】
また、本発明の窒化ケイ素質焼結体は、前記微細粒子は、シリコン成分の高い核と、前記焼結助剤成分の高い周辺部とからなることを特徴としている。すなわち、この微細粒子は少なくともＳｉ−Ｎ-Ｏ−Ｍｇ−ＲＥ組成を有しており、当該組成割合は核部分についてはシリコン成分が高く、かつ助剤成分として添加する（例えば、Ｍｇおよび希土類元素）成分量が小さい。一方周辺部分は、逆にシリコン成分が小さく、助剤成分量が多いという構成が望ましい。また当該微細粒子は、全体に非晶質相であることが望ましい。
【００１９】
前記酸化物換算含有量の合計が0.6ｗｔ％未満では焼結時の緻密化作用が不十分で相対密度が95％未満となり好ましくなく、10ｗｔ％超では窒化ケイ素質焼結体の第２のミクロ組織成分である熱伝導率の低い粒界相の量が過剰となり焼結体の熱伝導率が100Ｗ／（ｍ・K）未満になる。これら酸化物含有量の合計は0.6〜6ｗｔ％がより好ましい。尚且つ、（ＲＥｘＯｙ）／（ＭｇＯ）＞１であることが望ましく、この場合に特に高強度・高熱伝導性が向上する。これについては後述するが、希土類酸化物（ＲＥｘＯｙ）のイオン半径が酸化マグネシウム（ＭｇＯ）のイオン半径より大きく、窒化ケイ素粒子内に固溶するよりも析出した方が安定となることが新に知見されたことによる。また、本発明の窒化ケイ素質焼結体は、常温における熱伝導率が100〜300Ｗ／（ｍ・K）であり、常温における３点曲げ強度が600〜1500ＭＰａであり高強度・高熱伝導性に富んでいる。
【００２０】
また、本発明の窒化ケイ素質焼結体の製造方法は、β分率が30〜100％であり、酸素含有量が0.5ｗｔ％以下であり、平均粒子径が0.2〜10μmであり、アスペクト比が10以下である第一の窒化ケイ素質粉末１〜50重量部と、平均粒子径が0.2〜4μmであり、α型の第ニの窒化ケイ素粉末99〜50重量部と、希土類元素（ＲＥ）からなる群から選ばれた少なくとも１種の元素とＭｇとがその成分として含まれ、前記希土類元素を希土類酸化物（ＲＥ ｘ Ｏ ｙ ）に換算し、前記Ｍｇを酸化マグネシウム（ＭｇＯ）に換算したとき、これらを酸化物に換算した酸化物含有量の合計が 0.6 〜 10 ｗｔ％で、かつ（ＲＥ ｘ Ｏ ｙ ）／（ＭｇＯ）で表される重量比が１より大きい焼結助剤とを配合し、1800℃以上の焼結温度及び0.5ＭＰａ以上の窒素加圧雰囲気にて焼結する焼結工程を有することを特徴とする。ここで、前記焼結工程において、昇温時1400℃〜1600℃の温度で1〜10時間にわたる保持工程を少なくとも１回有し、かつこの保持工程における温度から前記焼結温度までの昇温速度を5.0℃/min以下にして、前記焼結工程を行なうことが好ましく、さらに好ましくは2.5℃/min以下である。
【００２１】
前記窒化ケイ素質粉末のβ分率が30％未満では成長核としての効果はあるものの部分的に核として作用するため、異常粒成長が起こり、最終的に得られる窒化ケイ素質焼結体のミクロ組織中に大きな粒子を均一分散できなくなり曲げ強度が低下する。したがって、窒化ケイ素質粉末のβ分率は30％以上が望ましい。
また前記窒化ケイ素質粉末の平均粒子径が0.2μm未満では前記同様に柱状粒子が均一に発達したミクロ組織を呈する窒化ケイ素質焼結体を得られず、熱伝導率および曲げ強度を高めることが困難である。また前記窒化ケイ素質粉末の平均粒子径が10μmより大きいと焼結体の窒化ケイ素質の緻密化が阻害される。したがって、窒化ケイ素質粉末の平均粒子径は0.2〜10μmが好ましい。
さらに、アスペクト比が10超の場合は窒化ケイ素質焼結体の緻密化が阻害され、結果として、常温における３点曲げ強度は600MPa未満になる。したがって、窒化ケイ素質粉末のアスペクト比を10以下とすることが好ましい。
【００２２】
前記焼結工程において、昇温時1400℃〜1600℃の温度で1〜10時間にわたる保持工程を入れること、およびこの保持温度から前記焼結温度までの昇温速度を5.0℃/min以下にすることは、焼結体の密度（焼結性）と最終ミクロ組織および窒化ケイ素粒子内への助剤成分および酸素成分の固溶量に影響を与える。すなわち1400℃〜1600℃の温度領域では、助剤成分とSi3N4粉末表面のSiO2成分が反応して液相を形成し、αからβへの相転位が起こり、続いて、粒成長が開始する。この温度領域で保持することにより成長核となるβ粒子の形状を均質化させる効果があり、この後の昇温工程における異常粒成長を抑制することができる。
また、助剤成分として、希土類酸化物とともにMg成分を添加する利点は、液相生成温度を低下させ、焼結性を改善できることにある。しかしながら、Mg成分は蒸気圧が高いため、焼結過程において焼結体の内部から表面部へのMg成分の拡散が進行する。このため、内部と表面部との組成差が生じ、とりわけ肉厚品を焼結する場合には両者間で色調差を呈し、さらには焼結体内部の密度ならびに強度が著しく低下するといった難点がある。この点1400℃〜1600℃の温度領域における保持工程を追加することで、この傾向を抑制する効果があり、緻密質かつ高強度の焼結体を得るために望ましい工程である。
【００２３】
次に、この保持温度から焼結温度への昇温速度を5.0℃/min以下とすると、Mg成分の急激な系外への揮発を抑制することができる。特に2.5℃/min以下とすると、最終焼結体中のMg量を制御することが容易となり、とりわけ、薄物シート焼結体に対して、各試料間でのMg量の組成差がなく、しいては、密度、強度等の諸特性において差が無くなり、製品歩留りならびに品質を安定させることができる。このため、焼結体中に気孔を生成させることなく、低熱伝導の粒界相を効率よく低減することができ、焼結体の熱伝導率向上に寄与する。また、溶解・再析出を繰り返す粒成長過程で、窒化ケイ素粒子内に取り込まれる助剤成分量ならびに酸素量を低減することができ、この効果も焼結体の熱伝導率向上に繋がる。したがって、昇温時の工程で保持すること、かつ保持温度から焼結温度までの昇温速度を5.0℃/min以下にすることは、焼結体の熱伝導率および強度を両立させるために望ましい工程である。
【００２４】
また、予め1650〜1850℃の焼結温度で成形体を予備焼成し、次いで1850〜1900℃の熱処理を行うと高熱伝導化が顕著になり120w/(m・K)を超える窒化ケイ素質焼結体を得られ特に好ましい。この熱処理による高熱伝導化は窒化ケイ粒子の成長と、蒸気圧の高いMgO基とした粒界相成分が効率よく窒化ケイ素質焼結体外へ揮発することの複合効果による。尚、1850℃〜1950℃の焼成温度にて、焼成時間を延長することで、上記同様の高熱伝導化の効果が達成できる。
【００２５】
【発明の実施の形態】
以下、本発明の実施形態について説明する。
本発明の窒化ケイ素質焼結体において、高温熱処理および焼成時間の延長により焼結体の熱伝導率は向上するが、これは、窒化ケイ素粒子の粒成長および焼結助剤成分の揮発による複合効果に加えて、窒化ケイ素粒子内に微細粒子が析出することが窒化ケイ素粒子自身の熱伝導率の上昇に影響を与えている。したがって、100 w/(m・K)以上の熱伝導率を得るためには、窒化ケイ素粒子内の微細粒子析出効果は有効である。更に、強度と熱伝導率を両立するためには、破壊の起点として作用する窒化ケイ素粒子の寸法を一定にし、この粒子内の高純度化作用を適用することが肝要である。
【００２６】
焼結助剤としてはＭｇおよびＹは有用であり、窒化ケイ素質原料粉末の緻密化に有効である。これらの元素は窒化ケイ素質焼結体を構成する第１ミクロ組織成分である窒化ケイ素質粒子に対する固溶度が小さいので、窒化ケイ素粒子、ひいては窒化ケイ素質焼結体の熱伝導率を高い水準に保つことができる。
また、Ｙと同様に窒化ケイ素質粒子に対する固溶度が小さく、焼結助剤として有用な元素として、Ｌａ，Ｃｅ，Ｐｒ，Ｎｄ，Ｐｍ，Ｓｍ，Ｅｕ，Ｇｄ，Ｔｂ,Ｄｙ，Ｈｏ，Ｅｒ，Ｔｍ，ＹｂおよびＬｕの群から選択される少なくとも１種の希土類元素が挙げられる。そのうち、温度および圧力が高くなり過ぎずに焼成ができる点でＬａ，Ｃｅ，Ｇｄ，ＤｙおよびＹｂの群から選択される少なくとも１種の希土類元素が好ましい。前記、微細粒子は、イオン半径の大きい元素が主で構成されており、焼結助剤として添加するＭｇを酸化マグネシウム（ＭｇＯ）換算し、また含有するＬａ，Ｙ,ＧｄおよびＹｂを含む希土類元素（ＲＥ）から選択される少なくとも１種の元素を酸化物（ＲＥｘＯｙ）換算した場合、ＲＥｘＯｙ／ＭｇＯ＞１である場合に、微細粒子が析出し易くなる。換言すれば焼結助剤として添加するＭｇＯ量が多い場合にこの微細粒子が析出し難たくなる。
【００２７】
この理由は、前述の様にこの微細粒子は、助剤成分とＳｉ、ＯおよびＮから構成されるが、Ｍｇ元素のイオン半径（Ｍｇ2+）半径：0.07nmは、窒化ケイ素（Si3N4）を構成するＳｉ元素のイオン半径：0.04nmに比較的近く、酸素と共に窒化ケイ素粒子内に固溶する形態が安定である。一方、希土類元素酸化物（ＲＥｘＯｙ）量が多い場合には、Ｙｂ以上の希土類元素のイオン半径（ＲＥx+）は、0.09nmでありＳｉ元素のイオン半径：0.04nmの２倍以上であり、またＭｇ元素のイオン半径（Ｍｇ2+）半径：0.07nmと比較して大きく、窒化ケイ素粒子内に固溶するよりも析出した形が安定となる。したがって、微細粒子を析出させるためには、焼結体の緻密化が達成できる範囲においては、希土類元素酸化物（ＲＥｘＯｙ）基であることが望ましい。ここで、微細粒子の粒子径が100nm超となると、それに伴い、窒化ケイ素粒子内に析出する100nm超の微細粒子の数が著しく増加する。微細粒子は、Si−N-O−Mg−Rからなるガラス相で構成されており、これ自身の熱伝導率は低い。このため、100nm超の微細粒子の存在が多くなると、逆に目的とする窒化ケイ素粒子自身の熱伝導率向上が達成できない。したがって、微細粒子は粒径100nm以下に制御することが肝要である。
【００２８】
窒化ケイ素焼結体の熱伝導率は、ミクロ組織と密接な関係にあり、これらを構成する窒化ケイ素粒子と粒界相の熱伝導率に支配される。後者は、主にガラス相として存在し、それらの熱伝導率は高々3W/(m・K)程度である。また、所定の熱処理あるいは、焼結後の冷却速度を緩やかにすることで粒界ガラス相を結晶化させた場合でも30W/(m・K)程度である。前者の熱伝導率は、理論値でAlNの319W/(m・K)に近い300W/(m・K)と推定されており、また実測値でも180W/(m・K)が得られている。したがって、焼結体の高熱伝導化は、窒化ケイ素粒子自身の熱伝導率が多く関与している。
ここで、窒化ケイ素粒子自身の熱伝導率を低下させる阻害要因として、粒内転位ならびに固溶元素がある。これらの阻害要因は、熱媒体であるフォノンの散乱を引き起こし熱伝達を著しく低減させる。このため、窒化ケイ素粒子の熱伝導率向上のため、しいては焼結体の熱伝導率向上のためには、これらの阻害因子を抑制することが肝要である。これら阻害要因のうち、粒内の固溶元素は焼結過程における液相生成段階でＳｉ、Ｎおよび助剤成分からなるSi−N-O−Mg−REを生成し、更に粒成長段階で比較的小さな粒子がこの液相に溶解して、続いてＳｉ、Ｎが大きな粒子の表面に再析出して粒成長が進行する。この際にＳｉ、Ｎに混じってＭｇ、ＲＥの助剤成分および酸素（O）も粒子表面に取り込まれる。上述した様に、元素のイオン半径が小さい程、この傾向は大きくなる。
【００２９】
よって、焼結後の最終のミクロ組織を構成する窒化ケイ素粒子内には、極微量の助剤成分および酸素が微細粒子に存在する。この固溶元素は例えばMg、Y、La、Gd、Yb等の希土類元素であり、これらを粒子内に微細に析出させれば、微細粒子の周りは高純度化され、粒子自身の熱伝導率は上昇する。このような固溶元素の存在が本発明の特徴的な点であり、これにより焼結体の高熱伝導化が達成できる。固溶元素の析出は上記した保持過程を含む焼結や焼結時間の延長また熱処理にて調整できる。しかしながら、焼結後の窒化ケイ素粒子内に固溶元素量が多い場合には、微細粒子析出による粒子の高純度化作用は起こらないため、適切な焼結助剤の選定ならびに焼結方法の適用が肝要である。
【００３０】
次に、窒化ケイ素焼結体の製造方法において、β分率が30〜100％の第一の窒化ケイ素質粉末と第二のα型窒化ケイ素質粉末との比率は１〜50ｗｔ％：99〜50ｗｔ％が好ましい。前記β分率が30〜100％の窒化ケイ素質粉末の比率が1ｗｔ％未満では成長核としての効果はあるものの、添加量が少ないために作用する成長核の数が少なく、異常粒成長が起こりミクロ組織中に大きな粒子を均一分散できなくなり、曲げ強度が低下する。また、50ｗｔ％超では成長核の数が多くなり、粒成長の過程で、粒子同士が互いに衝突するため成長阻害が起こり、強度は維持できるが、発達した柱状粒子からなる窒化ケイ素質焼結体のミクロ組織を得られず、従来に比べて高い熱伝導率を実現困難になる。
前記窒化ケイ素質粉末の酸素量を0.5ｗｔ％以下としたのは、前記窒化ケイ素質粉末を成長核として作用させて窒化ケイ素質焼結体を形成する場合、窒化ケイ素質焼結体を構成する窒化ケイ素質粒子内に固溶する酸素量は、成長核として用いる窒化ケイ素質粉末の酸素量に強く依存し、この窒化ケイ素質粉末の酸素量が高いほど窒化ケイ素質粒子内に固溶する酸素量が高くなる。そして窒化ケイ素質粒子中に含有される酸素により熱伝導媒体であるフォノンの散乱が発生し、窒化ケイ素質焼結体の熱伝導率が低下する。100W/(m・K)以上という従来の窒化ケイ素質焼結体では得られなかった高い熱伝導率を発現するには、窒化ケイ素質粉末の含有酸素量を0.5wt%以下に抑えて、最終的に得られる窒化ケイ素質焼結体の酸素量を低減することが必要不可欠である。
【００３１】
窒化ケイ素質粉末中のＦｅ含有量およびＡｌ含有量がそれぞれ100ppm超では窒化ケイ素粒子内にＦｅまたはＡｌが顕著に固溶し、この固溶部分で熱伝導媒体であるフォノンの散乱を生じ、窒化ケイ素質焼結体の熱伝導率を低下させる。したがって100W/m.K以上の熱伝導率を得るには窒化ケイ素質粉末中のＦｅ含有量およびＡｌ含有量をそれぞれ100ppm以下に制御することが肝要である。
【００３２】
本発明の窒化ケイ素質焼結体からなる基板は高強度、高靭性ならびに高熱伝導率の特性を生かして、パワ−半導体用基板またはマルチチップモジュ−ル用基板などの各種基板、あるいはペルチェ素子用熱伝板、または各種発熱素子用ヒ−トシンクなどの電子部品用部材に好適である。
例えば窒化ケイ素質焼結体を半導体素子用基板として用いた場合、半導体素子の作動に伴う繰り返しの熱サイクルを受けたときの基板のクラックの発生が抑えられ、耐熱衝撃性ならびに耐熱サイクル性が著しく向上し、信頼性に優れたものとなる。また、高出力化および高集積化を指向する半導体素子を搭載した場合でも、熱抵抗特性の劣化が少なく、優れた放熱特性を発揮する。さらに、優れた機械的特性により本来の基板材料としての機能だけでなく、それ自体が構造部材を兼ねることができるため、基板ユニット自体の構造を簡略化できる。
【００３３】
また、あるいは本発明の窒化ケイ素質焼結体をペルチェ素子用熱伝板として用いた場合、ペルチェ素子の印加電圧の極性の入れ替えに伴う繰り返し熱サイクルを受けたときの前記基板のクラックの発生が抑えられ、耐熱サイクル性が著しく向上し、信頼性に優れたものとなる。また、ゼーベック素子熱伝板として用いる場合、吸熱側では６００℃前後の高温になるため、ここでも耐熱サイクル性かつ耐熱衝撃性が要求されるが、これに本発明の窒化ケイ素質焼結体を用いた場合には、これらの寿命特性が大幅に向上し、信頼性の優れたものとなる。
【００３４】
また、本発明の窒化ケイ素質焼結体は、上述の電子部品用部材以外に熱衝撃および熱疲労の耐熱抵抗特性が要求される材料に幅広く利用できる。構造用部材として、各種の熱交換器部品や熱機関用部品、アルミニウムや亜鉛等の金属溶解の分野で用いられるヒーターチューブ、ストークス、ダイカストスリーブ、溶湯攪拌用プロペラ、ラドル、あるいは熱電対保護管等に適用できる。また、アルミニウム、亜鉛等の溶融金属めっきラインで用いられるシンクロール、サポートロール、軸受、あるいは軸等に適用することにより、急激な加熱や冷却に対して耐割れ性に富んだ部材となり得る。また、鉄鋼あるいは非鉄の加工分野では、圧延ロール、スキーズロール、ガイドローラ、線引きダイス、あるいは工具用チップ等に用いれば、被加工物との接触時の放熱性が良好なため、耐熱疲労性および耐熱衝撃性を改善することができ、これにより摩耗が少なく、熱応力割れを生じにくくできる。
【００３５】
さらに、スパッタターゲット部材にも適用でき、例えば磁気記録装置のＭＲヘッド、ＧＭＲヘッド、またはＴＭＲヘッドなどに用いられる電気絶縁膜の形成や、熱転写プリンターのサーマルヘッドなどに用いられる耐摩耗性皮膜の形成に好適である。スパッタして得られる被膜は、本質的に高熱伝導特性を持つとともに、スパッタレートも十分高くでき、被膜の電気的絶縁耐圧が高いものとなる。このため、このスパッタターゲットで形成したＭＲヘッド、ＧＭＲヘッド、またはＴＭＲヘッド用の電気絶縁性被膜は高熱伝導ならびに高耐電圧の特性を有するので、素子の高発熱密度化や絶縁性被膜の薄膜化が図れる。また、このスパッタターゲットで形成したサ−マルヘッド用の耐摩耗性被膜は、窒化ケイ素本来の特性により耐摩耗性が良好であることはもとより、高熱伝導性のため熱抵抗が小さくできるので印字速度を高めることができる。
【００３６】
以下、実施例により本発明を説明するが、それら実施例により本発明が限定されるものではない。
（実施例１）
β化率が30％以上の第―の窒化ケイ素質粉末１〜50wt%と、平均粒径が0.7〜1.2μｍ、酸素量が0.5〜2.0wt%のα型の第二の窒化ケイ素質粉末を、1.0wt%または2.0wt%のMgOと、3wt%または6.0wt%のGd2O3あるいは表１に示す焼結助剤を添加した混合粉末を作製した。なお、第ニの窒化ケイ素粉末の割合は、第―の窒化ケイ素粉末と焼結助剤粉末のバランスとした。さらに2wt%の分散剤（商品名：レオガ-ドGP）を配合し、エタノールを満たしたボ−ルミル容器中に投入し、次いで混合した。得られた混合物を真空乾燥し、次いで目開き150μmの篩を通して造粒した。次に、プレス機により直径20mm×厚さ10mmおよび直径100mm×厚さ15mmのディスク状の成形体を圧力3tonのCIP成形により得た。次いで1850℃〜1950℃、0.7〜0.9MPa（7〜9気圧）の窒素ガス雰囲気中で5〜40時間焼成した。なお、焼結工程において、昇温時1400℃〜1600℃の温度で1〜10時間にわたる保持工程を設け、かつこの保持温度から前記焼結温度までの昇温速度を5.0℃/min以下にした。個々の試料の製造条件は表１の試料No.1〜15の欄に示す。
【００３７】
また、得られた窒化ケイ素質焼結体の窒化ケイ素粒子内の微細粒子の観察は、透過型電子顕微鏡（日立製作所製ＨＦ２０００）にて観察倍率×10,000倍から600,000倍で行った。更に、微細粒子の組成分析は付属のエネルギー分散型分析装置にて評価した。図１〜図４は、本発明の窒化ケイ素焼結体（表１中のNo.１からNo.４の試料）のＴＥＭ観察像の写真である。また、図５は比較例のＴＥＭ観察像（表３中のNo.３１の試料）の写真である。図６〜図９は、微細粒子の高分解能観察像（表１中のNo.１からNo.４の試料）の写真、図１０（表１中のNo.１試料）及び図１１（表１中のNo.２試料）は、微細粒子の核および周辺部のＳＴＥＭ観察像の写真である。
【００３８】
次に得られた窒化ケイ素質焼結体から、直径10mm×厚さ3mmの熱伝導率および密度測定用の試験片、ならびに縦3mm×横4mm×長さ40mmの曲げ試験片を採取した。密度はマイクロメ−タにより寸法を測定し、また重量を測定し、算出した。熱伝導率はレーザーフラッシュ法により常温での比熱および熱拡散率を測定し熱伝導率を算出した。３点曲げ強度は常温にてＪＩＳ Ｒ１６０６に準拠して測定を行った。
以上の製造条件の概略および評価結果を、表１および表２の試料No.１〜15に示す。
【００３９】
（比較例１）
表１記載の試料No.31〜42の製造条件とした以外は実施例１と同様にして評価した。以上の製造条件の概略および評価結果を、表１および表２の試料No.31〜42に示す。
【００４０】
【表１】
【表２】
表１および表２に示したように、窒化ケイ素粒子内に微細粒子が認めれた焼結体については、いずれも100W/(m・K)以上の熱伝導率と600MPa以上の曲げ強度が得られた。しかも微細粒子の存在割合が増すほど熱伝導率が向上する傾向が確認できた。微細粒子が認められた焼結体について用いた焼結助剤成分のREｘOｙ／MgO比は１以上であった。一方、窒化ケイ素粒子内に微細粒子が認められない焼結体については、いずれも100W/(m・K)未満の熱伝導率となった。これに加えて、試料No.37〜42については、焼結工程において昇温時1400℃〜1600℃の温度で保持しない場合、あるいはこの保持温度から焼結温度までの昇温速度を5.0℃/min超とした場合には、熱伝導率および強度ともに著しく低下した。
【００４３】
図１〜図５のTEM観察像、図６〜図９の微細粒子の高分解能電子顕微鏡（HREM）観察像及び図１０、図１１の微細粒子の核および周辺部のSTEM観察像について考察する。
図１〜図４に、焼結助剤としてGd2O3（図１）、Yb2O3（図２）、Y2O3（図３）およびLa2O3（図４）を用いた本発明例の透過型電子顕微鏡（TEM）像を示す。また、図５に比較例のTEM観察像を示す。
図１〜図４より、いずれにおいても窒化ケイ素粒子内に微細粒子が存在する。図１では右下部に8〜45nmの範囲で点在、図２では中央右部に10〜60nm範囲で点在、図３では中央左右部に8〜60nm範囲nmで点在、図４では右上部に4〜85nmの範囲で点在していることが分かり、これらの微細粒子の粒径はいずれも100nm以下であった。一方、図５の比較例については、この様な微細粒子は観察されなかった。なお、別の観察視野においても確認されなかった。ここで、微細粒子の粒子径が100nm超となると、それに伴い、窒化ケイ素粒子内に析出する100nm超の微細粒子の数が著しく増加し、所望の窒化ケイ素粒子自身の熱伝導率向上に寄与しない。
【００４４】
次に図６〜図９に、焼結助剤としてGd2O3（図６）、Yb2O3（図７）、Y2O3（図８）およびLa2O3（図９）を用いた本発明例の高分解能観察（HREM）像を示す。図６〜図９はそれぞれ図１〜図４で観察された微細粒子についての観察像である。図６〜図９から、窒化ケイ素粒子内に析出する微細粒子は、ランダムな格子像および電子回折像がガラス相特有のハローパターンを示したことから非晶質相からなることが判明した。更に、図７のHREM像においては、組成のことなる核７と周辺部８からなることを確認した。TEM−EDX分析の結果から、核はSi成分が高く、MgおよびRE成分（本発明ではYbが該当）が低く、また周辺部の組成は、これと逆の評価結果であった。なお、HREM像においては極微小領域を長時間観察した場合には、電子線によるダメージのため核と周辺部を分離して観察することは困難であるが、このHREM像は微細粒子の構成要素の分離、さらに組成の定量化まで言及できた点で非常に優れている。
図１０および図１１は、焼結助剤としてGd2O3（図１０）およびYb2O3（図１１）を用いた本発明例の走査透過型電子顕微鏡（STEM）像を示す。これらの図は、それぞれ、図１および図２にて観察された微細粒子についてのSTEM像である。STEM像は、ナノレベルの微小領域を観察する場合、特に、組成や成分量の僅かな差を画像コントラストとして表現するのに有効な観察方法である。図１０および図１１に示した様に、個々の微細粒子は、核と周辺部から構成されることが確認でき、核はSi成分が高く、かつMgおよびRE（本発明ではGd、Ybが該当）が低く、一方、周辺部はこれとは逆の組成であることが判明した。
【００４５】
（実施例２）
β化率が30％以上、酸素含有量が0.5wt%以下、平均粒子径が1μm〜10μm、アスペクト比が10以下の第―の窒化ケイ素質粉末を１〜50ｗｔ％と平均粒径が0.7〜1.2μｍ、酸素量が0.5〜2.0wt%のα型の第二の窒化ケイ素質粉末に1wt%のMgO、3wt%％Gd2O3の焼結助剤を添加した混合粉末を作製した。次いで、アミン系の分散剤を2ｗｔ％添加したトルエン・ブタノール溶液を満たしたボールミルの樹脂製ポット中に作製した混合粉末および粉砕媒体の窒化ケイ素製ボールを投入し、48時間湿式混合した。次いで、前記ポット中の混合粉末100重量部に対しポリビニル系の有機バインダーを15重量部および可塑剤（ジメチルフタレ−ト）を5重量部添加し、次いで48時間湿式混合しシート成形用スラリーを得た。この成形用スラリーを調整後、ドクターブレード法によりグリーンシート成形した。次いで、成形したグリーンシートを空気中400〜600℃で2〜5時間加熱することにより、予め添加し有機バインダー成分を十分に脱脂（除去）した。次いで脱脂体を0.9MPa（9気圧）の窒素雰囲気中で1900℃×10時間の焼成を行い、その後室温に冷却した。焼結工程においては、昇温時1400℃〜1600℃の温度で1〜10時間にわたる保持工程を設け、かつこの保持温度から前記焼結温度までの昇温速度を2.0℃/minとした。得られた窒化ケイ素質焼結体シートに機械加工を施し縦50mm×横50mm×厚さ0.6mmの半導体モジュール用の基板を製造した。
【００４６】
この窒化ケイ素質焼結体製基板を用いて図１２に示す回路基板を作製した。図１２において、回路基板１１は作製した前記縦50mm×横50mm×厚さ0.6mmの寸法の窒化ケイ素質焼結体製基板１２の表面に銅製回路板１３を設け、前記基板１２の裏面に銅板１４をろう材１５により接合して構成されている。
この回路基板１１に対し、３点曲げ強度の評価および耐熱サイクル試験を行った。その結果、曲げ強度が600MPa以上と大きく、回路基板１１の実装工程における締め付け割れおよびはんだ付け工程時の熱応力に起因するクラックの発生する頻度がほぼ見られなくなり、回路基板を使用した半導体装置の製造歩留まりを大幅に改善できることが実証された。また、耐熱サイクル試験は、−40℃での冷却を20分、室温での保持を10分および180℃における加熱を20分とする昇温/降温サイクルを１サイクルとし、これを繰り返し付与し、基板部にクラック等が発生するまでのサイクル数を測定した。その結果、1000サイクル経過後においても窒化ケイ素質焼結体製基板１２の割れや銅製回路板１３の剥離はなく、優れた耐久性と信頼性を兼備することが確認された。また、1000サイクル経過後においても耐電圧特性の低下は発生しなかった。
【００４７】
最後に実施例１および２で用いたβ分率が30％以上の窒化ケイ素粉末について述べておく、含有酸素量がＳｉＯ２換算で2.0ｗｔ％未満、平均粒子径0.2〜2.0μmのイミド分解法による窒化ケイ素質粉末をＢＮ製るつぼに充填し、次いで常圧〜1.0MPa（10気圧）のＮ2雰囲気中にて1400℃〜1950℃で１〜20時間加熱する熱処理を施し、次いで室温まで冷却した。得られた窒化ケイ素質粉末のβ分率は90〜100％であり、酸素含有量は0.2〜0.4ｗｔ％であった。図１３に得られた窒化ケイ素質粉末例のＳＥＭ観察像を示す。当該粉末のβ分率は100％、酸素量は0.2ｗｔ％、ＦｅおよびＡｌ量はそれぞれ、50ppmおよび40ppmであった。当該粉末には粒子の長軸方向と平行に溝部が形成されており、これは気相を介して粒成長が起こる場合の特徴で、特に酸素量が微量であるほど顕著となる。
得られた窒化ケイ素質粉末のＦｅ、Ａｌの不純物分析はプラズマ発光分析（ICP）法により行った。また、酸素含有量は赤外線加熱吸収法により測定した。
【００４８】
また得られた窒化ケイ素質粉末のβ分率はＣｕ―Ｋα線を用いたＸ線回折強度比から式(1)により求めた。
β分率（％）＝ {(Iβ (101)+Iβ (210))/(Iβ (101)+Iβ (210)+Iα (102)+Iα (201))}×100 (1)
Iβ（１０１） ：β型Ｓｉ３Ｎ４の(101)面回折ピ-ク強度，
Iβ（２１０） ：β型Ｓｉ３Ｎ４の(210)面回折ピ-ク強度，
Iα（１０２） ：α型Ｓｉ３Ｎ４の(102)面回折ピ-ク強度，
Iα（２１０） ：α型Ｓｉ３Ｎ４の(210)面回折ピ-ク強度。
【００４９】
また、得られた窒化ケイ素質粉末の平均粒子径および平均アスペクト比は、SEM観察にて観察倍率×2000倍で得られたSEM写真を用い、200μm×500μm視野面積内にある計500個の窒化ケイ素質粒子を無作為に選定して画像解析装置により最小径と最大径を測定し、その平均値を求めて評価した。
得られた窒化ケイ素質粉末は、β分率が30％以上、平均粒子径が0.5〜10μm、アスペクト比が10以下、ＦｅおよびＡｌの含有量が、いずれも100ppm以下、また、酸素含有量は、0.5wt%以下であった。
【００５０】
【発明の効果】
以上の通り、本発明の窒化ケイ素質焼結体は、窒化ケイ素粒子内にMgあるいはY、La、Gd、Yb等の希土類元素の内の少なくとも１種の元素と、酸素元素とを含む粒径100nm以下の微細粒子の存在により、本来有する高強度/高靭性に加えて高い熱伝導率を具備したものとなる。これは、高温・高圧焼結といったコストの高い焼成法、焼成装置を必要とせずに製造することが出来る。また、これを半導体素子用基板として用いた場合に半導体素子の作動に伴う繰り返しの熱サイクルによって基板にクラックが発生することが少なく、耐熱衝撃性ならびに耐熱サイクル性を著しく向上することができる。
【図面の簡単な説明】
【図１】本発明例の希土類酸化物にGd2O3を用いた場合の窒化ケイ焼結体の透過型電子顕微鏡（ＴＥＭ）観察写真を示す。
【図２】本発明例の希土類酸化物にYb2O3を用いた場合の窒化ケイ焼結体の透過型電子顕微鏡（ＴＥＭ）観察写真を示す。
【図３】本発明例の希土類酸化物にY2O3を用いた場合の窒化ケイ焼結体の透過型電子顕微鏡（ＴＥＭ）観察写真を示す。
【図４】本発明例の希土類酸化物にLa2O3を用いた場合の窒化ケイ焼結体の透過型電子顕微鏡（ＴＥＭ）観察写真を示す。
【図５】比較例の窒化ケイ焼結体の透過型電子顕微鏡（ＴＥＭ）観察写真を示す。
【図６】本発明例の希土類酸化物にGd2O3を用いた場合の窒化ケイ焼結体において、窒化ケイ素粒子内に析出した微細粒子の高分解能観察写真（ＨＲＥＭ）を示す。
【図７】本発明例の希土類酸化物にYb2O3を用いた場合の窒化ケイ焼結体において、窒化ケイ素粒子内に析出した微細粒子の高分解能観察写真（ＨＲＥＭ）を示す。
【図８】本発明例の希土類酸化物にY2O3を用いた場合の窒化ケイ焼結体において、窒化ケイ素粒子内に析出した微細粒子の高分解能観察写真（ＨＲＥＭ）を示す。
【図９】本発明例の希土類酸化物にLa2O3を用いた場合の窒化ケイ焼結体において、窒化ケイ素粒子内に析出した微細粒子の高分解能観察写真（ＨＲＥＭ）を示す。
【図１０】本発明例の希土類酸化物にGd2O3を用いた場合の窒化ケイ焼結体において、窒化ケイ素粒子内に析出した微細粒子の走査型透過電子顕微鏡写真（ＳＴＥＭ）を示す。
【図１１】本発明例の希土類酸化物にYb2O3を用いた場合の窒化ケイ焼結体において、窒化ケイ素粒子内に析出した微細粒子の走査型透過電子顕微鏡写真（ＳＴＥＭ）を示す。
【図１２】本発明例の窒化ケイ素質焼結体を用いた回路基板の要部断面図を示す。
【図１３】本発明例の窒化ケイ素焼結体の製造に用いた窒化ケイ素質粉末のSEM観察像写真を示す。
【符号の説明】
１１：回路基板
１２：基板
１３：銅製回路板
１４：銅板
１５：ろう材。[0001]
BACKGROUND OF THE INVENTION
The present invention provides a high strength and high heat suitable as a member for electronic parts such as a semiconductor substrate and a heat sink for a heating element, or a structural member such as a member for general machinery, a member for molten metal, or a member for a heat engine. The present invention relates to a silicon nitride sintered body rich in conductivity, a method for producing the same, and a circuit board configured using the silicon nitride sintered body.
[0002]
[Prior art]
Silicon nitride sintered bodies have excellent heat resistance, low thermal expansion, thermal shock resistance, and corrosion resistance against metals in addition to mechanical properties such as high-temperature strength and wear resistance. It is used for various structural members such as members, engine members, steelmaking machine members, or molten metal resistant members. In addition, it is used as an electrical insulating material by utilizing high insulating properties.
[0003]
2. Description of the Related Art In recent years, with the development of semiconductor devices that generate large amounts of heat, such as high-frequency transistors and power ICs, there is an increasing demand for ceramic substrates having high thermal conductivity in order to obtain good heat dissipation characteristics in addition to electrical insulation. As such a ceramic substrate, an aluminum nitride substrate is used, but mechanical strength, fracture toughness, etc. are low, and there is a problem that cracking occurs due to tightening in the assembly process of the substrate unit. In addition, in a circuit board in which an Si semiconductor element is mounted on an aluminum nitride substrate, the thermal expansion difference between Si and the aluminum nitride substrate is large. There's a problem.
[0004]
Accordingly, although the thermal conductivity is inferior to that of an aluminum nitride substrate, a substrate made of a highly thermally conductive silicon nitride sintered body having a thermal expansion coefficient close to that of Si and excellent in mechanical strength, fracture toughness and heat fatigue resistance has attracted attention. Proposals have been made.
[0005]
For example, Japanese Patent Laid-Open No. 4-175268 discloses that Al and oxygen, which are substantially made of silicon nitride and both contained as impurities, are not more than 3.5% by weight, and the density is not less than 3.15 Mg / m 3 (3.15 g / cm 3). A silicon nitride sintered body having a thermal conductivity of 40 W / (m · K) or more is described.
[0006]
Japanese Patent Application Laid-Open No. 9-30866 discloses that 85-99% by weight of β-type silicon nitride grains and the balance are composed of oxide or oxynitride grain boundary phases, and Mg, Ca, Containing 0.5 to 10% by weight of at least one element selected from Sr, Ba, Y, La, Ce, Pr, Nd, Sm, Gd, Dy, Ho, Er, and Yb, and Al in the grain boundary phase Silicon nitride-based firing having an element content of 1% by weight or less, a porosity of 5% or less, and a proportion of β-type silicon nitride grains having a minor axis diameter of 5 μm or more of 10-60% by volume Conjunctions are described.
[0007]
In addition, in Japanese Ceramic Society 1996 Annual Conference Proceedings 1G11 and 1G12 and JP-A-10-194842, columnar silicon nitride particles or whiskers are added to the raw material powder in advance, and the doctor blade method or extrusion molding method is used. A silicon nitride sintered body having a heat conductivity in a specific direction increased by imparting anisotropy to the heat conduction by forming a sintered body in which the particles are two-dimensionally oriented and sintering. Is described.
[0008]
JP-A-2001-19557 discloses thermal conductivity by controlling the total amount of oxygen, Al, Ca, and Fe impurities in silicon nitride grains to 1500 ppm or less and the minor axis diameter to 2 μm or more. And a silicon nitride sintered body with improved mechanical properties and a method for producing the same.
[0009]
Furthermore, in JP-A-2002-29848, columnar whiskers are added in advance to the raw material powder, and in the firing process, a microstructure in which grains are selectively grown using the whiskers as nuclei is constructed, thereby increasing the thermal conductivity. An improved silicon nitride sintered body and a method for producing the whisker are described.
[0010]
Japan Ceramic Society 1998 Lecture Proceedings 2B04 includes a compact of silicon nitride powder sintered in 1.0MPa nitrogen gas at 2000 ℃ x 4hr, and then in 2MP at 30MPa nitrogen gas at 2200 ℃ x 4hr. It describes that a silicon nitride sintered body having a high thermal conductivity of 100 w / (m · K) or more can be produced by performing a heat treatment at a high temperature and a high pressure. It is described that the development of high thermal conductivity is related to the growth of silicon nitride particles in the sintered body and the hexagonal precipitate phase in the silicon nitride particles by high-temperature heat treatment.
[0011]
[Problems to be solved by the invention]
Hereinafter, the above-described problems of the prior art will be described in order.
First, in Japanese Patent Laid-Open No. 4-175268, a thermal conductivity of 40 W / (m · K) or more has been obtained, but recently, a material having further improved thermal conductivity and excellent mechanical strength is desired. .
Further, in the methods described in JP-A-9-30866, JP-A-10-194842, etc., seed crystals or whiskers as growth nuclei are obtained in order to obtain huge columnar particles in the silicon nitride sintered body. Is added in advance, and firing in a nitrogen atmosphere at 2000 ° C. or higher and 10.1 MPa (100 atm) or higher is essential. Therefore, special high-temperature / high-pressure equipment such as hot press or HIP is required, resulting in cost increase. Further, since the molding process for obtaining a molded body in which silicon nitride particles are oriented is complicated, there is a problem that productivity is significantly reduced.
[0012]
Moreover, the silicon nitride sintered body described in JP-A-2001-19557 improves the thermal conductivity of the particles themselves by reducing the amount of impurities in the silicon nitride grains (denoted as a purification effect), This is characterized by improving the thermal conductivity of the sintered body. Further, Zr and / or Hf is selected as an additive useful for promoting the purification effect, and 0.5 wt% to 3.0 wt% are added as oxides. However, when oxides of Zr and Hf are added, ZrN and HfN having electrical conductivity are generated in the grain boundary phase by easily reacting with N in Si3N4 during the sintering process. Therefore, there is a problem that the electrical insulation that is essentially required for a ceramic substrate cannot be maintained and it is difficult to use as an insulating substrate for a power semiconductor module that operates at a high frequency.
[0013]
Further, in the production method described in JP-A-2002-29848, the average circularity of the silicon nitride particles to be formed is 0.8 or more, the β conversion is 10% or more and less than 80%, and the oxygen amount is 0.5 to 1.8 mass. %, Specific surface area is 12-22m2One or more selected from the group consisting of rare earth oxides, silicon oxides, and magnesium oxides are added to the silicon nitride powder that is / g, so that the total is 2.5 to 14% by mass. -Is added in an amount of 0.1 to 8.5% by mass, mixed, molded, and sintered in a nitriding atmosphere. Further, the silicon nitride whisker is hydrothermally treated in advance at a temperature higher than the boiling point of water (for example, at a temperature range of 110 ° C. to 140 ° C.) in order to improve sinterability. However, the hydrothermal treatment of the silicon nitride whiskers used in this example oxidizes the surface and acts as an auxiliary2By increasing the components, the sinterability can be improved, but on the other hand, as described in the examples, hydrothermal treatment to obtain a dense sintered body requires treatment at 120 ° C. for 96 hours and the process is complicated. become. In addition, although the sinterability can be improved by hydrothermal treatment, the oxygen content of the sintered body, that is, the oxygen content in the silicon nitride particles cannot be reduced, and it is difficult to obtain a high thermal conductive material. .
[0014]
Next, the sintered body described in the 2nd B04 of the 1998 Annual Meeting of the Ceramic Society of Japan was fired at 2000 ° C in 1MPa nitrogen gas, followed by high-temperature and high-pressure heat treatment at 2200 ° C in 30MPa nitrogen gas. Therefore, there is an advantage that a high thermal conductivity of 100 W / (m · K) or more can be obtained. Furthermore, it is explained that the mechanism of high thermal conductivity is that the hexagonal precipitate phase in the silicon nitride particles is involved by the high temperature heat treatment in addition to the growth of the silicon nitride particles in the sintered body. In other words, the auxiliary component composed of Y-Nd-Si-O is incorporated into the silicon nitride particles during sintering and grain growth and becomes a solid solution, and the Y-Nd-Si-O composition during heat treatment and cooling at high temperature As an amorphous phase, it precipitates in the silicon nitride particles, and a part of the precipitate is considered to be crystallized. From the above, in order to obtain the above sintered body, a special high-temperature / high-pressure facility is required, resulting in an increase in cost. Further, since the heat treatment is applied after sintering, there is a problem that the productivity is remarkably lowered. Further, a detailed composition analysis and observation have not been made on the precipitated phase in the silicon nitride particles in the sintered body, and the relevance to the improvement in thermal conductivity has not been clarified.
[0015]
The present invention has been made in view of the above-mentioned conventional problems, and does not require an expensive baking method such as high-temperature and high-pressure baking at 2000 ° C. or higher and 10.1 MPa (100 atm) or higher, and has excellent mechanical strength. Another object of the present invention is to provide a high thermal conductivity type silicon nitride sintered body having a higher thermal conductivity than the prior art without anisotropy in the direction of thermal conduction.
Another object of the present invention is to provide a high thermal conductivity type silicon nitride sintered body having an increased thermal conductivity by examining in detail the composition and form of fine particles precipitated in the silicon nitride particles.
In addition, the present invention provides high thermal conductivity by defining the β fraction of silicon nitride powder, the amount of oxygen contained, the amount of impurities, the mixing ratio with α-type silicon nitride powder and the sintering process including the holding process. An object of the present invention is to provide a silicon nitride sintered body having high strength and a method for producing the same.
Another object of the present invention is to provide a circuit board with good heat dissipation constructed using the above-described silicon nitride sintered body having high strength and high thermal conductivity.
[0016]
[Means for Solving the Problems]
In order to achieve the above-mentioned problems, the present inventors consciously precipitate fine particles containing at least oxygen and a sintering aid component in the silicon nitride particles, thereby improving the thermal conductivity of the silicon nitride particles themselves. It was found that a silicon nitride sintered body having a thermal conductivity of 100 W / (m · K) or more and sufficient bending strength can be obtained stably. At this time, the sintering aid component is MgO group, so that the sinterability is improved and MgO and (RExOy) Was found to be effective to contain with a specific amount and a specific ratio. Further, in the method for producing a silicon nitride sintered body, the sintering process includes the powder characteristics such as the β fraction of the silicon nitride powder to be used, the oxygen content, the mixing ratio of impurities and α powder, and the holding process. It was found that it is important to specify the above. Thus, the present invention has been achieved.
[0017]
That is, the silicon nitride sintered body of the present invention isA silicon nitride-based sintered body in which at least one element selected from the group consisting of rare earth elements (RE) and Mg are added as components of a sintering aid, the rare earth elements being rare earth oxides (RE). x O y ) And when the Mg is converted to magnesium oxide (MgO), the total oxide content converted to these oxides is 0.6 ~ Ten wt% and (RE x O y ) / (MgO) is a weight ratio greater than 1 and contains the sintering aid component, oxygen, nitrogen, and silicon, and the average particle size 100nm The following fine particles are precipitated in the silicon nitride particles constituting the silicon nitride sintered body.It is characterized by that. The silicon nitride sintered body of the present invention ismagnification 10,000 In transmission electron micrographs more than double,In silicon nitride particlesAverage particle sizeFine particles of 100nm or less5 / μ m 2 PrecipitationIt is characterized by that.The fine particles include a nucleus having a high silicon component and a peripheral portion having a high sintering aid component. The silicon nitride sintered body of the present invention is characterized in that the fine particles are in an amorphous phase.The fine particles are formed by re-precipitating the auxiliary components incorporated into the silicon nitride particles with the growth of the silicon nitride particles during the firing process, and the high thermal conductivity of the silicon nitride particles themselves. Contributes to At this time, in the observation image of direct magnification of 10,000 times or more with a transmission electron microscope (TEM), the fine particles having an average particle diameter of 100 nm or less are 5 / μm in the silicon nitride particles.2The presence of the above is desirable, and the thermal conductivity of the sintered body is improved by the precipitation phenomenon and ratio of the fine particles.
[0018]
In addition, the silicon nitride sintered body of the present invention includes the fine particles.Consists of a core with a high silicon component and a peripheral part with a high sintering aid componentIt is characterized by that.That is,The fine particles have at least a Si—N—O—Mg—RE composition, and the composition ratio is about the core portion.siliconThe amount of the component is high and added as an auxiliary component (for example, Mg and rare earth elements)small. On the other hand, the surrounding areasiliconA configuration in which the component is small and the amount of the auxiliary component is large is desirable. The fine particles are desirably in an amorphous phase as a whole.
[0019]
If the total oxide content is less than 0.6 wt%, the densification effect during sintering is insufficient and the relative density is less than 95%, which is not preferable, and if it exceeds 10 wt%, the second micro of the silicon nitride sintered body is not preferable. The amount of the grain boundary phase having a low thermal conductivity, which is a structural component, becomes excessive, and the thermal conductivity of the sintered body becomes less than 100 W / (m · K). The total of these oxide contents is more preferably 0.6-6 wt%. And (RExOy) / (MgO)> 1, and in this case, particularly high strength and high thermal conductivity are improved. As will be described later, rare earth oxides (RExOy) Is larger than the ionic radius of magnesium oxide (MgO), and it has been newly found that precipitation is more stable than solid solution in silicon nitride particles. The silicon nitride sintered body of the present invention has a thermal conductivity of 100 to 300 W / (m · K) at room temperature, a three-point bending strength of 600 to 1500 MPa at room temperature, and has high strength and high thermal conductivity. Rich.
[0020]
The method for producing a silicon nitride sintered body of the present invention has a β fraction of 30 to 100%, an oxygen content of 0.5 wt% or less, an average particle size of 0.2 to 10 μm, and an
[0021]
If the β fraction of the silicon nitride powder is less than 30%, it has an effect as a growth nucleus, but partially acts as a nucleus. Large particles cannot be uniformly dispersed in the tissue, and the bending strength decreases. Therefore, the β fraction of the silicon nitride powder is desirably 30% or more.
In addition, when the average particle diameter of the silicon nitride powder is less than 0.2 μm, it is impossible to obtain a silicon nitride sintered body having a microstructure in which columnar particles are uniformly developed, as described above, and it is possible to increase thermal conductivity and bending strength. Have difficulty. If the average particle size of the silicon nitride powder is larger than 10 μm, densification of the silicon nitride material in the sintered body is hindered. Accordingly, the average particle size of the silicon nitride powder is preferably 0.2 to 10 μm.
Further, when the aspect ratio is more than 10, densification of the silicon nitride sintered body is inhibited, and as a result, the three-point bending strength at room temperature is less than 600 MPa. Therefore, the aspect ratio of the silicon nitride powder is preferably 10 or less.
[0022]
In the sintering step, a holding step for 1 to 10 hours at a temperature of 1400 ° C. to 1600 ° C. at the time of temperature rise is added, and the heating rate from the holding temperature to the sintering temperature is 5.0 ° C./min or less. This affects the density (sinterability) of the sintered body, the final microstructure, and the solid solution amount of the auxiliary component and oxygen component in the silicon nitride particles. That is, in the temperature range of 1400 ° C. to 1600 ° C., the auxiliary component reacts with the
The advantage of adding the Mg component together with the rare earth oxide as an auxiliary component is that the liquid phase formation temperature can be lowered and the sinterability can be improved. However, since the Mg component has a high vapor pressure, the Mg component diffuses from the inside of the sintered body to the surface portion during the sintering process. For this reason, there is a difference in composition between the inside and the surface portion, particularly when a thick product is sintered, there is a difference in color tone between the two, and the density and strength inside the sintered body are significantly reduced. is there. By adding a holding step in the temperature range of 1400 ° C. to 1600 ° C. in this respect, there is an effect of suppressing this tendency, which is a desirable step for obtaining a dense and high-strength sintered body.
[0023]
Next, when the rate of temperature increase from the holding temperature to the sintering temperature is 5.0 ° C./min or less, rapid volatilization of the Mg component outside the system can be suppressed. In particular, when the temperature is 2.5 ° C./min or less, it becomes easy to control the Mg amount in the final sintered body, and in particular, there is no compositional difference in Mg amount between the samples for the thin sheet sintered body. Therefore, there is no difference in various characteristics such as density and strength, and the product yield and quality can be stabilized. For this reason, it is possible to efficiently reduce the grain boundary phase having low thermal conductivity without generating pores in the sintered body, which contributes to the improvement of the thermal conductivity of the sintered body. In addition, in the grain growth process in which dissolution and reprecipitation are repeated, the amount of the auxiliary component and the amount of oxygen taken into the silicon nitride particles can be reduced, and this effect also leads to an improvement in the thermal conductivity of the sintered body. Therefore, it is desirable to maintain the temperature rising step from the holding temperature to the sintering temperature and to keep the temperature rising rate to 5.0 ° C./min or less in order to achieve both the thermal conductivity and strength of the sintered body. It is a process.
[0024]
In addition, pre-firing the compact at a sintering temperature of 1650 to 1850 ° C in advance, followed by heat treatment at 1850 to 1900 ° C, a significant increase in thermal conductivity becomes significant and silicon nitride sintering exceeding 120 w / (m · K) A body is obtained, which is particularly preferable. The high thermal conductivity by this heat treatment is due to the combined effects of the growth of silicon nitride particles and the efficient vaporization of MgO-based grain boundary phase components out of the silicon nitride sintered body. In addition, by extending the firing time at a firing temperature of 1850 ° C. to 1950 ° C., the same effect of increasing thermal conductivity can be achieved.
[0025]
DETAILED DESCRIPTION OF THE INVENTION
Hereinafter, embodiments of the present invention will be described.
In the silicon nitride sintered body of the present invention, the thermal conductivity of the sintered body is improved by high-temperature heat treatment and extension of the firing time. This is a combination of grain growth of silicon nitride particles and volatilization of the sintering aid component. In addition to the effect, the precipitation of fine particles in the silicon nitride particles affects the increase in the thermal conductivity of the silicon nitride particles themselves. Therefore, in order to obtain a thermal conductivity of 100 w / (m · K) or more, the fine particle precipitation effect in the silicon nitride particles is effective. Furthermore, in order to achieve both strength and thermal conductivity, it is important to keep the size of silicon nitride particles acting as a starting point of fracture constant and to apply a high-purity action in the particles.
[0026]
Mg and Y are useful as sintering aids and are effective in densifying the silicon nitride material powder. Since these elements have a low solid solubility with respect to the silicon nitride particles that are the first microstructure components constituting the silicon nitride sintered body, the thermal conductivity of the silicon nitride particles, and hence the silicon nitride sintered body, is high. Can be kept in.
Further, like Y, the solubility in silicon nitride particles is small, and elements useful as a sintering aid are La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Dy, Ho, Er. , Tm, Yb, and Lu, and at least one rare earth element. Of these, at least one rare earth element selected from the group of La, Ce, Gd, Dy, and Yb is preferable in that firing can be performed without excessively increasing the temperature and pressure. The fine particles are mainly composed of an element having a large ionic radius. Mg added as a sintering aid is converted into magnesium oxide (MgO), and the rare earth element containing La, Y, Gd, and Yb is contained. At least one element selected from (RE) is an oxide (RExOy) When converted, RExOyWhen / MgO> 1, fine particles are likely to precipitate. In other words, when the amount of MgO added as a sintering aid is large, it becomes difficult to precipitate these fine particles.
[0027]
The reason for this is that, as described above, this fine particle is composed of an auxiliary component and Si, O and N, but the ionic radius of Mg element (Mg2+) Radius: 0.07nm is relatively close to the ionic radius of Si element constituting silicon nitride (Si3N4): 0.04nm, and the form of solid solution in silicon nitride particles with oxygen is stable. On the other hand, rare earth element oxides (RExOy) When the amount is large, the ion radius (RE) of rare earth elements equal to or greater than Ybx +) Is 0.09 nm, and the ionic radius of Si element is more than twice that of 0.04 nm, and the ionic radius of Mg element (Mg)2+) Radius: Larger than 0.07 nm, and the precipitated form is more stable than solid solution in silicon nitride particles. Therefore, in order to precipitate fine particles, rare earth element oxides (RE) are used as long as densification of the sintered body can be achieved.xOy) Group. Here, when the particle diameter of the fine particles exceeds 100 nm, the number of fine particles of over 100 nm precipitated in the silicon nitride particles increases accordingly. The fine particles are composed of a glass phase composed of Si—N—O—Mg—R, and its own thermal conductivity is low. For this reason, if the presence of fine particles exceeding 100 nm increases, the intended improvement in the thermal conductivity of the silicon nitride particles themselves cannot be achieved. Therefore, it is important to control the fine particles to have a particle size of 100 nm or less.
[0028]
The thermal conductivity of the silicon nitride sintered body is closely related to the microstructure, and is governed by the thermal conductivity of the silicon nitride particles constituting these and the grain boundary phase. The latter exists mainly as a glass phase, and their thermal conductivity is at most about 3 W / (m · K). Further, even when the grain boundary glass phase is crystallized by a predetermined heat treatment or a slow cooling rate after sintering, it is about 30 W / (m · K). The thermal conductivity of the former is estimated to be 300 W / (m · K), which is close to AlN's 319 W / (m · K) as a theoretical value, and 180 W / (m · K) is also obtained as an actual measurement. . Therefore, high thermal conductivity of the sintered body is largely related to the thermal conductivity of the silicon nitride particles themselves.
Here, there are an intragranular dislocation and a solid solution element as an inhibiting factor that lowers the thermal conductivity of the silicon nitride particles themselves. These inhibiting factors cause scattering of the phonon, which is a heat medium, and significantly reduce heat transfer. For this reason, in order to improve the thermal conductivity of the silicon nitride particles and thus to improve the thermal conductivity of the sintered body, it is important to suppress these inhibitory factors. Among these inhibiting factors, the solid solution elements in the grains produce Si-NO-Mg-RE consisting of Si, N and auxiliary components in the liquid phase formation stage in the sintering process, and are relatively small in the grain growth stage. Particles dissolve in this liquid phase, and then Si and N reprecipitate on the surface of the large particles, and grain growth proceeds. At this time, the auxiliary components of Mg and RE and oxygen (O) are also incorporated into the particle surface in a mixture with Si and N. As described above, this tendency increases as the ionic radius of the element decreases.
[0029]
Therefore, in the silicon nitride particles constituting the final microstructure after sintering, a very small amount of auxiliary components and oxygen are present in the fine particles. This solid solution element is, for example, a rare earth element such as Mg, Y, La, Gd, Yb, etc. If these are finely precipitated in the particles, the periphery of the fine particles is highly purified, and the thermal conductivity of the particles themselves Will rise. The presence of such a solid solution element is a characteristic point of the present invention, and this makes it possible to achieve high thermal conductivity of the sintered body. The precipitation of the solid solution element can be adjusted by sintering including the above-described holding process, by extending the sintering time, or by heat treatment. However, when there is a large amount of solid solution elements in the sintered silicon nitride particles, the high-purity action of the particles due to the precipitation of fine particles does not occur, so the selection of an appropriate sintering aid and the application of the sintering method Is essential.
[0030]
Next, in the method for producing a silicon nitride sintered body, the ratio of the first silicon nitride powder having a β fraction of 30 to 100% and the second α-type silicon nitride powder is 1 to 50 wt%: 99 to 50 wt% is preferred. If the ratio of the silicon nitride powder having a β fraction of 30 to 100% is less than 1 wt%, there is an effect as a growth nucleus, but since the addition amount is small, the number of acting growth nuclei is small and abnormal grain growth occurs. Large particles cannot be uniformly dispersed in the microstructure, and the bending strength decreases. On the other hand, if it exceeds 50 wt%, the number of growth nuclei increases, and the grains collide with each other in the process of grain growth, resulting in growth inhibition and maintaining strength, but a silicon nitride sintered body composed of developed columnar particles. It is difficult to achieve a high thermal conductivity as compared with the prior art.
The reason why the oxygen content of the silicon nitride powder is 0.5 wt% or less is that when the silicon nitride powder is used as a growth nucleus to form a silicon nitride sintered body, the silicon nitride sintered body is formed. The amount of oxygen dissolved in the silicon nitride particles strongly depends on the amount of oxygen in the silicon nitride powder used as the growth nucleus. The higher the amount of oxygen in the silicon nitride powder, the more oxygen dissolved in the silicon nitride particles. The amount becomes higher. And the scattering of the phonon which is a heat conductive medium generate | occur | produces with the oxygen contained in a silicon nitride particle, and the heat conductivity of a silicon nitride sintered compact falls. In order to express the high thermal conductivity of 100 W / (mK) or more, which was not possible with conventional silicon nitride sintered bodies, the content of oxygen in the silicon nitride powder was suppressed to 0.5 wt% or less, and the final It is indispensable to reduce the amount of oxygen in the silicon nitride-based sintered body that is obtained.
[0031]
When the Fe content and Al content in the silicon nitride powder are each over 100 ppm, Fe or Al is remarkably dissolved in the silicon nitride particles, and the phonon, which is a heat conduction medium, is scattered in the solid solution portion. The thermal conductivity of the silicon sintered body is lowered. Therefore, in order to obtain a thermal conductivity of 100 W / m.K or more, it is important to control the Fe content and Al content in the silicon nitride powder to 100 ppm or less, respectively.
[0032]
The substrate comprising the silicon nitride sintered body of the present invention makes use of the characteristics of high strength, high toughness and high thermal conductivity, and is used for various substrates such as a power semiconductor substrate or a multichip module substrate, or for a Peltier device. It is suitable for a member for electronic parts such as a heat transfer plate or a heat sink for various heating elements.
For example, when a silicon nitride-based sintered body is used as a substrate for a semiconductor element, the generation of cracks in the substrate when subjected to repeated thermal cycles accompanying the operation of the semiconductor element is suppressed, and the thermal shock resistance and thermal cycle resistance are remarkably high. Improves reliability. Further, even when a semiconductor element oriented to higher output and higher integration is mounted, the thermal resistance characteristics are hardly deteriorated and excellent heat dissipation characteristics are exhibited. Furthermore, the structure of the substrate unit itself can be simplified because not only functions as the original substrate material but also the structural member itself can be obtained due to excellent mechanical characteristics.
[0033]
Or, when the silicon nitride sintered body of the present invention is used as a heat transfer plate for a Peltier element, the occurrence of cracks in the substrate when subjected to repeated thermal cycles accompanying the change of polarity of the applied voltage of the Peltier element It is suppressed, the heat cycle resistance is remarkably improved, and the reliability is excellent. In addition, when used as a Seebeck element heat transfer plate, the heat absorption side has a high temperature of about 600 ° C., so heat cycle resistance and thermal shock resistance are also required here. When used, these life characteristics are greatly improved and the reliability becomes excellent.
[0034]
Moreover, the silicon nitride sintered body of the present invention can be widely used for materials that require heat resistance characteristics of thermal shock and thermal fatigue in addition to the above-described electronic component member. As structural members, various heat exchanger parts, heat engine parts, heater tubes, Stokes, die casting sleeves, molten metal stirring propellers, ladles, thermocouple protection tubes, etc. used in the field of melting metals such as aluminum and zinc Applicable to. Moreover, by applying to sink rolls, support rolls, bearings, shafts, etc. used in molten metal plating lines of aluminum, zinc, etc., it can be a member having excellent crack resistance against rapid heating and cooling. Also, in the steel or non-ferrous processing field, if it is used for rolling rolls, squeeze rolls, guide rollers, wire drawing dies, tool tips, etc., it has good heat dissipation when it comes into contact with the workpiece, so it is resistant to heat fatigue. In addition, the thermal shock resistance can be improved, thereby reducing wear and making it difficult to cause thermal stress cracking.
[0035]
Furthermore, it can be applied to sputter target members, for example, the formation of electrical insulation films used for MR heads, GMR heads, or TMR heads of magnetic recording devices, and the formation of wear-resistant films used for thermal heads of thermal transfer printers. It is suitable for. The film obtained by sputtering inherently has high heat conduction characteristics, a sufficiently high sputtering rate, and high electrical withstand voltage of the film. For this reason, the electrical insulating film for MR head, GMR head, or TMR head formed with this sputter target has characteristics of high thermal conductivity and high withstand voltage, so that the heat generation density of the element is increased and the insulating film is made thinner. Can be planned. In addition, the thermal-resistant coating for thermal heads formed with this sputter target not only has good wear resistance due to the inherent characteristics of silicon nitride, but also has high thermal conductivity, so the thermal resistance can be reduced, so the printing speed can be reduced. Can be increased.
[0036]
EXAMPLES Hereinafter, although an Example demonstrates this invention, this invention is not limited by these Examples.
(Example 1)
a first silicon nitride powder of 1 to 50 wt% with a β conversion rate of 30% or more, an α-type second silicon nitride powder with an average particle size of 0.7 to 1.2 μm and an oxygen content of 0.5 to 2.0 wt%. 1.0wt% or 2.0wt% MgO and 3wt% or 6.0wt% Gd2OThreeOr the mixed powder which added the sintering auxiliary agent shown in Table 1 was produced. The ratio of the second silicon nitride powder was the balance between the first silicon nitride powder and the sintering aid powder. Further, 2 wt% of a dispersant (trade name: Leocard GP) was blended, put into a ball mill container filled with ethanol, and then mixed. The obtained mixture was vacuum-dried and then granulated through a sieve having an opening of 150 μm. Next, a disk-shaped molded body having a diameter of 20 mm × thickness of 10 mm and a diameter of 100 mm × thickness of 15 mm was obtained by CIP molding at a pressure of 3 tons. Next, it was calcined in a nitrogen gas atmosphere at 1850 ° C. to 1950 ° C. and 0.7 to 0.9 MPa (7 to 9 atm) for 5 to 40 hours. In the sintering process, a holding process was performed for 1 to 10 hours at a temperature of 1400 ° C. to 1600 ° C. at the time of temperature increase, and the rate of temperature increase from this holding temperature to the sintering temperature was 5.0 ° C./min or less. . The manufacturing conditions of the individual samples are shown in the columns of Sample Nos. 1 to 15 in Table 1.
[0037]
Further, observation of fine particles in the silicon nitride particles of the obtained silicon nitride sintered body was performed with a transmission electron microscope (HF2000 manufactured by Hitachi, Ltd.) at an observation magnification of 10,000 to 600,000 times. Furthermore, the composition analysis of the fine particles was evaluated with an attached energy dispersion type analyzer. 1 to 4 are photographs of TEM observation images of the silicon nitride sintered body of the present invention (samples No. 1 to No. 4 in Table 1). FIG. 5 is a photograph of a TEM observation image of the comparative example (No. 31 sample in Table 3). 6 to 9 are photographs of high-resolution observation images of fine particles (samples No. 1 to No. 4 in Table 1), FIG. 10 (No. 1 sample in Table 1) and FIG. 11 (Table 1). No. 2 sample) is a photograph of the STEM observation image of the core and peripheral part of the fine particles.
[0038]
Next, a test piece for measuring thermal conductivity and density having a diameter of 10 mm × thickness of 3 mm and a bending test piece having a length of 3 mm × width of 4 mm × length of 40 mm were collected from the obtained silicon nitride-based sintered body. The density was calculated by measuring the dimensions with a micrometer and measuring the weight. The thermal conductivity was calculated by measuring the specific heat and thermal diffusivity at room temperature by the laser flash method. The three-point bending strength was measured at room temperature according to JIS R1606.
The outline and evaluation results of the above production conditions are shown in Sample Nos. 1 to 15 in Tables 1 and 2.
[0039]
(Comparative Example 1)
Evaluation was performed in the same manner as in Example 1 except that the production conditions of Sample Nos. 31 to 42 shown in Table 1 were used. The outline and evaluation results of the above production conditions are shown in Sample Nos. 31 to 42 in Tables 1 and 2.
[0040]
[Table 1]
[Table 2]
As shown in Tables 1 and 2, all sintered bodies with fine particles found in silicon nitride particles have a thermal conductivity of 100 W / (m · K) or more and a bending strength of 600 MPa or more. It was. Moreover, it was confirmed that the thermal conductivity tends to improve as the proportion of fine particles increases. The RExOy / MgO ratio of the sintering aid component used for the sintered body in which fine particles were observed was 1 or more. On the other hand, all of the sintered bodies in which fine particles were not observed in the silicon nitride particles had a thermal conductivity of less than 100 W / (m · K). In addition to this, for sample Nos. 37 to 42, the temperature increase rate from the holding temperature to the sintering temperature was 5.0 ° C / In the case of exceeding min, both the thermal conductivity and strength were significantly reduced.
[0043]
The TEM observation image of FIGS. 1-5, the high-resolution electron microscope (HREM) observation image of the fine particle of FIGS. 6-9, and the STEM observation image of the nucleus and peripheral part of the fine particle of FIG. 10, FIG.
1 to 4 show transmission electron microscope (TEM) images of examples of the present invention using Gd2O3 (FIG. 1), Yb2O3 (FIG. 2), Y2O3 (FIG. 3) and La2O3 (FIG. 4) as sintering aids. Indicates. FIG. 5 shows a TEM observation image of the comparative example.
1 to 4, in any case, fine particles are present in the silicon nitride particles. 1 is dotted in the range of 8 to 45 nm in the lower right, FIG. 2 is dotted in the range of 10 to 60 nm in the center right, FIG. 3 is dotted in the range of 8 to 60 nm in the center left and right, and in FIG. It was found that the particles were scattered in the range of 4 to 85 nm, and the particle size of these fine particles was 100 nm or less. On the other hand, such fine particles were not observed in the comparative example of FIG. It was not confirmed in another observation field. Here, when the particle size of the fine particles exceeds 100 nm, the number of fine particles of over 100 nm deposited in the silicon nitride particles increases remarkably and does not contribute to the improvement of the thermal conductivity of the desired silicon nitride particles themselves. .
[0044]
Next, FIGS. 6 to 9 show high resolution observation (HREM) of the present invention example using Gd2O3 (FIG. 6), Yb2O3 (FIG. 7), Y2O3 (FIG. 8) and La2O3 (FIG. 9) as sintering aids. Show the image. 6 to 9 are observation images of the fine particles observed in FIGS. From FIG. 6 to FIG. 9, it was found that the fine particles precipitated in the silicon nitride particles consist of an amorphous phase because random lattice images and electron diffraction images showed a halo pattern peculiar to the glass phase. Further, in the HREM image of FIG. 7, it was confirmed that the composition was composed of a
10 and 11 show scanning transmission electron microscope (STEM) images of examples of the present invention using Gd2O3 (FIG. 10) and Yb2O3 (FIG. 11) as sintering aids. These figures are STEM images of the fine particles observed in FIGS. 1 and 2, respectively. The STEM image is an effective observation method for expressing a slight difference in composition and component amount as an image contrast, particularly when observing a nano-level minute region. As shown in FIG. 10 and FIG. 11, it can be confirmed that each fine particle is composed of a nucleus and a peripheral part, the nucleus has a high Si component, and Mg and RE (in the present invention, Gd and Yb are applicable). ) Was low, while the periphery was found to have the opposite composition.
[0045]
(Example 2)
1 to 50 wt% of the first silicon nitride powder having a β conversion rate of 30% or more, an oxygen content of 0.5 wt% or less, an average particle diameter of 1 μm to 10 μm, and an aspect ratio of 10 or less, and an average particle diameter of 0.7 to 1.2μm, 0.5% to 2.0wt% α-type second silicon nitride powder with 1wt% MgO, 3wt %% Gd2OThreeA mixed powder added with the sintering aid was prepared. Next, the mixed powder prepared in a resin pot of a ball mill filled with a toluene / butanol solution to which 2 wt% of an amine-based dispersant was added and a silicon nitride ball as a grinding medium were charged and wet mixed for 48 hours. Next, 15 parts by weight of a polyvinyl organic binder and 5 parts by weight of a plasticizer (dimethyl phthalate) were added to 100 parts by weight of the mixed powder in the pot, and then wet-mixed for 48 hours to obtain a sheet forming slurry. . After the molding slurry was adjusted, green sheets were molded by the doctor blade method. Next, the formed green sheet was heated in the air at 400 to 600 ° C. for 2 to 5 hours to add in advance and sufficiently degrease (remove) the organic binder component. Next, the degreased body was baked at 1900 ° C. for 10 hours in a nitrogen atmosphere of 0.9 MPa (9 atm), and then cooled to room temperature. In the sintering step, a holding step for 1 to 10 hours was provided at a temperature of 1400 ° C. to 1600 ° C. at the time of temperature increase, and the rate of temperature increase from this holding temperature to the sintering temperature was 2.0 ° C./min. The obtained silicon nitride sintered sheet was machined to produce a substrate for a semiconductor module having a length of 50 mm, a width of 50 mm, and a thickness of 0.6 mm.
[0046]
A circuit board shown in FIG. 12 was produced using this silicon nitride sintered body substrate. In FIG. 12, a circuit board 11 is provided with a copper circuit board 13 on the surface of a silicon nitride-based sintered board 12 having dimensions of 50 mm in length, 50 mm in width, and 0.6 mm in thickness. 14 is joined by a
The circuit board 11 was subjected to three-point bending strength evaluation and heat cycle test. As a result, the bending strength is as high as 600 MPa or more, and the frequency of occurrence of cracking due to the tightening crack in the mounting process of the circuit board 11 and the thermal stress during the soldering process is almost not seen, and the semiconductor device using the circuit board is It has been demonstrated that manufacturing yield can be significantly improved. In addition, the heat cycle test is a cycle of temperature increase / decrease with 20 minutes of cooling at −40 ° C., 10 minutes of holding at room temperature and 20 minutes of heating at 180 ° C. The number of cycles until a crack or the like occurred in the substrate portion was measured. As a result, even after 1000 cycles, the silicon nitride sintered substrate 12 was not cracked and the copper circuit board 13 was not peeled, and it was confirmed that both excellent durability and reliability were obtained. Moreover, the withstand voltage characteristics did not deteriorate even after 1000 cycles.
[0047]
Finally, silicon nitride powder having a β fraction of 30% or more used in Examples 1 and 2 will be described.2A silicon nitride powder obtained by imide decomposition having an average particle size of less than 2.0 wt% and an average particle size of 0.2 to 2.0 μm is filled into a BN crucible, and then N to 1.0 MPa (10 atm) under normal pressure.2Heat treatment was performed by heating at 1400 ° C. to 1950 ° C. for 1 to 20 hours in an atmosphere, and then cooled to room temperature. The obtained silicon nitride powder had a β fraction of 90 to 100% and an oxygen content of 0.2 to 0.4 wt%. FIG. 13 shows an SEM observation image of the obtained silicon nitride powder example. The β fraction of the powder was 100%, the oxygen content was 0.2 wt%, and the Fe and Al contents were 50 ppm and 40 ppm, respectively. Grooves are formed in the powder in parallel with the major axis direction of the particles, and this is a feature when grain growth occurs through the gas phase, and becomes more remarkable as the amount of oxygen is very small.
Impurity analysis of Fe and Al of the obtained silicon nitride powder was performed by plasma emission analysis (ICP) method. The oxygen content was measured by an infrared heat absorption method.
[0048]
Further, the β fraction of the obtained silicon nitride powder was obtained from the X-ray diffraction intensity ratio using Cu—Kα ray by the formula (1).
β fraction (%) = {(Iβ (101)+ Iβ (210)) / (Iβ (101)+ Iβ (210)+ Iα (102)+ Iα (201))} × 100 (1)
Iβ (101) : Β-type Si3N4(101) plane diffraction peak intensity,
Iβ (210) : Β-type Si3N4Of (210) plane diffraction peak intensity,
Iα (102) : Α-type Si3N4(102) plane diffraction peak intensity,
Iα (210) : Α-type Si3N4(210) Diffraction peak intensity.
[0049]
In addition, the average particle diameter and average aspect ratio of the obtained silicon nitride-based powder are SEM photographs obtained at an observation magnification of × 2000 magnification by SEM observation, and a total of 500 nitrides within a 200 μm × 500 μm visual field area. Silicon particles were selected at random, and the minimum and maximum diameters were measured by an image analyzer, and the average values were obtained and evaluated.
The obtained silicon nitride powder has a β fraction of 30% or more, an average particle size of 0.5 to 10 μm, an aspect ratio of 10 or less, and Fe and Al contents of 100 ppm or less, and an oxygen content of 0.5 wt% or less.
[0050]
【The invention's effect】
As described above, the silicon nitride-based sintered body of the present invention has a particle size that includes at least one element of rare earth elements such as Mg, Y, La, Gd, and Yb and oxygen element in the silicon nitride particles. The presence of fine particles of 100 nm or less provides high thermal conductivity in addition to the inherent high strength / toughness. This can be manufactured without requiring a high-cost baking method such as high temperature and high pressure sintering and a baking apparatus. Further, when this is used as a substrate for a semiconductor element, cracks are hardly generated in the substrate due to repeated thermal cycles accompanying the operation of the semiconductor element, and the thermal shock resistance and the thermal cycle resistance can be remarkably improved.
[Brief description of the drawings]
1 shows a transmission electron microscope (TEM) observation photograph of a silicon nitride sintered body when Gd 2
FIG. 2 shows a transmission electron microscope (TEM) observation photograph of a sintered silicon nitride when Yb2O3 is used for the rare earth oxide of the present invention.
FIG. 3 shows a transmission electron microscope (TEM) observation photograph of a silicon nitride sintered body when Y 2
FIG. 4 shows a transmission electron microscope (TEM) observation photograph of a silicon nitride sintered body when La2O3 is used as the rare earth oxide of the present invention.
FIG. 5 shows a transmission electron microscope (TEM) observation photograph of a silicon nitride sintered body of a comparative example.
FIG. 6 shows a high-resolution observation photograph (HREM) of fine particles precipitated in silicon nitride particles in a sintered silicon nitride when Gd 2
FIG. 7 shows a high-resolution observation photograph (HREM) of fine particles precipitated in silicon nitride particles in a sintered silicon nitride when Yb2O3 is used for the rare earth oxide of the present invention.
FIG. 8 shows a high-resolution observation photograph (HREM) of fine particles precipitated in silicon nitride particles in a silicon nitride sintered body when Y 2
FIG. 9 shows a high-resolution observation photograph (HREM) of fine particles precipitated in silicon nitride particles in a sintered silicon nitride when La2O3 is used for the rare earth oxide of the present invention.
FIG. 10 shows a scanning transmission electron micrograph (STEM) of fine particles precipitated in silicon nitride particles in a sintered silicon nitride when Gd 2
FIG. 11 shows a scanning transmission electron micrograph (STEM) of fine particles precipitated in silicon nitride particles in a sintered silicon nitride when Yb2O3 is used for the rare earth oxide of the present invention.
FIG. 12 is a cross-sectional view of a principal part of a circuit board using a silicon nitride sintered body according to an example of the present invention.
FIG. 13 shows a SEM observation image photograph of the silicon nitride powder used in the production of the silicon nitride sintered body of the present invention example.
[Explanation of symbols]
11: Circuit board
12: Substrate
13: Copper circuit board
14: Copper plate
15: Brazing material.
Claims (1)
β分率が30〜100％であり、酸素含有量が0.5ｗｔ％以下であり、平均粒子径が0.2〜10μmであり、アスペクト比が10以下である第一の窒化ケイ素質粉末１〜50重量部と、
平均粒子径が0.2〜4μmであり、α型の第ニの窒化ケイ素粉末99〜50重量部と、
希土類元素（ＲＥ）からなる群から選ばれた少なくとも１種の元素とＭｇとがその成分として含まれ、前記希土類元素を希土類酸化物（ＲＥｘＯｙ）に換算し、前記Ｍｇを酸化マグネシウム（ＭｇＯ）に換算したとき、これらを酸化物に換算した酸化物含有量の合計が0.6〜10ｗｔ％で、かつ（ＲＥｘＯｙ）／（ＭｇＯ）で表される重量比が１より大きい焼結助剤とを配合し、
1800℃以上の焼結温度及び0.5ＭＰａ以上の窒素加圧雰囲気にて焼結する焼結工程を有し、
該焼結工程において、昇温時1400℃〜1600℃の温度で1〜10時間にわたる保持工程を少なくとも１回有し、この保持工程における温度から前記焼結温度までの昇温速度を5.0℃/min以下にして、前記焼結工程を行なうことを特徴とする窒化ケイ素質焼結体の製造方法。A method for producing a silicon nitride sintered body rich in high strength and high thermal conductivity,
1 to 50 weights of first silicon nitride powder having a β fraction of 30 to 100%, an oxygen content of 0.5 wt% or less, an average particle diameter of 0.2 to 10 μm, and an aspect ratio of 10 or less And
The average particle diameter is 0.2 to 4 μm, α type second silicon nitride powder 99 to 50 parts by weight,
At least one element selected from the group consisting of rare earth elements (RE) and Mg are included as its components, the rare earth elements are converted into rare earth oxides (RE x O y ), and the Mg is converted into magnesium oxide ( When converted to MgO), the total oxide content converted to oxide is 0.6 to 10 wt%, and the weight ratio expressed by (RE x O y ) / (MgO) is larger than 1. With the aid,
Have a sintering step of sintering at 1800 ° C. or more sintering temperature and 0.5MPa or more nitrogen pressurized atmosphere,
In the sintering step, has at least one holding step over 1 to 10 hours at a temperature of in heating 1400 ° C. to 1600 ° C., the heating rate from the temperature in the holding step to the sintering temperature 5.0 ° C. / The method for producing a silicon nitride sintered body, wherein the sintering step is performed at a minimum value or less.
Priority Applications (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP2002121345A JP3775335B2 (en) | 2002-04-23 | 2002-04-23 | Silicon nitride sintered body, method for producing silicon nitride sintered body, and circuit board using the same |
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| JP4556162B2 (en) * | 2004-03-11 | 2010-10-06 | 日立金属株式会社 | Silicon nitride-based sintered body, method for producing the same, and circuit board using the same |
| DE112006000992B4 (en) | 2005-04-28 | 2014-03-27 | Hitachi Metals, Ltd. | Silicon nitride substrate, silicon nitride substrate manufacturing method, silicon nitride circuit board using the silicon nitride substrate, and semiconductor module using the same |
| JP4478198B2 (en) | 2007-03-22 | 2010-06-09 | 日本特殊陶業株式会社 | Inserts and cutting tools |
| KR101582704B1 (en) | 2008-07-03 | 2016-01-05 | 히타치 긴조쿠 가부시키가이샤 | Silicon nitride board, method for manufacturing the silicon nitride board, and silicon nitride circuit board and semiconductor module using the silicon nitride board |
| JP5142889B2 (en) * | 2008-08-27 | 2013-02-13 | 京セラ株式会社 | Silicon nitride sintered body, manufacturing method thereof, circuit board, and power semiconductor module |
| JPWO2013121861A1 (en) * | 2012-02-13 | 2015-05-11 | 三井金属鉱業株式会社 | Silicon nitride sintered body and method for producing the same |
| WO2013146789A1 (en) * | 2012-03-26 | 2013-10-03 | 日立金属株式会社 | Sintered silicon nitride substrate and process for producing same |
| US11865624B2 (en) | 2018-08-28 | 2024-01-09 | Kyocera Corporation | Insert and cutting tool |
| WO2021117829A1 (en) * | 2019-12-11 | 2021-06-17 | 宇部興産株式会社 | Plate-like silicon nitride-based sintered body and method for producing same |
| WO2024111402A1 (en) * | 2022-11-21 | 2024-05-30 | 株式会社 東芝 | Silicon nitride sintered compact, silicon nitride substrate, silicon nitride circuit board, and semiconductor device |
| WO2024177400A1 (en) * | 2023-02-24 | 2024-08-29 | 주식회사 아모센스 | Composition for manufacturing silicon nitride substrate and silicon nitride substrate manufactured thereby |
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